ournal J Am Cerum Soc 85 [101 3457-61( Control of Composition and Structure in Laminated Silicon Nitride/Boron Nitride Composites Chang-an Wang, Yong Huang, Qingfeng Zan, Linhua Zou, and Shengyou Cai State Key Laboratory of New Ceramics and Fine Processing, Department of Materials Science and Engineering Tsinghua University. Beijing, 100084, People's Republic of China Based on a biomimetic design, Si, N,/BN composites with MPam. Thereafter, various methods including tape casting and laminated structures have been prepared and investigated slip casting have been applied to make ceramic-matrix composites through composition control and structure design. To further with multilayered structures. The laminated design is considered to improve the mechanical properties of the composites, SiN be a promising method to overcome the brittle nature of ceramics matrix layers were reinforced by SiC whiskers and BN ser pa- However, the bending strength of the laminated materials at room rating layers were modified by adding Si,N, or Al2O3. The temperature is considerably reduced because of the presence of results showed that the addition of Sic whiskers in the SiaNa weak separating layers. Laminated composites consist of matrix layers could greatly improve the apparent fracture ceramic-matrix and interfacial separating layers. The former toughness(reaching 28.1 MPam"), at the same time keeping mainly determines the bulk strength and the latter gives rise to the the higher bending strength (reaching 651.5 MPa) of the high toughness due to the deflection of crack. Therefore, to composites. Additions of 50 wt% AL,O, or 10 wt% Si,N, to BN improve the mechanical properties of these composites, two ways interfacial layers had a beneficial effect on the strength and may be used: adjusting the composition and structure of the toughness of the laminated Si, N,/BN composites. Through interfacial separating layers and/or reinforcing the ceramic matrix observation of microstructure by SEM. multilevel toughening mechanisms contributing to high toughness of the laminated In the current paper, laminated Si, N,/BN composites were Si, N/BN composites were present as the first-level toughening prepared to imitate the structure of mollusk shell: thin Si, N, layers mechanisms from BN interfacial lavers as crack deflection were laminated like the aragonite plates with BN separating layer bifurcation, and pull-out of matrix sheets, and the secondary as the organic joining. To improve the mechanial properties of toughening mechanism from whiskers in matrix layers. laminated Si, N, /BN composites, SiC whiskers were added to Si,N, matrix layers for further reinforcement of matrix layers, and AL, O, or Si N, was incorporated into the BN layers for the I. Introduction purpose of adjusting the bonding strength between Si N, layers S LICON NITRIDE is known as an important high-temperature structural ceramic with excellent mechanical properties from IL. Experimental Procedure ambient to elevated temperatures, but the brittleness limits its applications. Therefore, various approaches have been chosen to Si,N,(Founder High-Tech Ceramic Corp, China) ration of particles and whiskers. - etc but improvement in the AL O,(purity>99.9%) were ball-milled with 20 wt% Sic toughness of the material has been only limited. However, it is well known that some natural biomaterials, such to achieve a homogeneous mixture. The slurry was dried and bamboo, wood, and nacre, have specific composite structures sieved through a 60-mesh screen; then the powder was mixed with and exhibit good mechanical properties in that they have high organic binders and then rolled into a sheet about 200 um thick by toughness and excellent damage resistance, compared with artifi- rolling compaction. The green sheets were subsequently coated formed by aragonite layers about I um thick joined by a kind of at 450C for 2 h in air. the green body was hot-pressed in a mortar of proteins. This particular configuration imparts over 1 graphite resistance furnace in flowing N2 at 1820 C for 1.5 hFor order of magnitude higher bending strength and toughness than comparison, Si, N, monolithic samples were fabricated using the those of aragonite single crystals. These natural composites may same composition and processing of Si, N,/BN laminated compos- give us some insights into making better structural materials with high toughness through biomimetic design The mechanical properties of the laminated Si, N,/BN compos Recently, some researchers-> introduced laminated or multi ites were determined on an A-2000 Shimadzu universal material layered structures to improve the mechanical properties of struc testing machine. Six samples were tested for each property. The tural ceramics. In 1990, Clegg reported preparing a laminated Sic loads were applied perpendicular to the plane of laminates composite separated by graphite layers, and the apparent fracture Bending strength was determined by three-point bend testing (test toughness of the composite was improved from 3.6 to 17.7 crosshead speed of 0.5 mm/min. The tensile surface of the samples was polished with diamond paste down to I um, and the corners were chamfered with a 15 mm diamond grinding wheel N. Claussen--contributing editor As for the fracture toughness of laminated ceramics, even though there are only a very few references mentioned, in which the fracture toughness was measured by SENB. SEVNB, and the indentation method, up to the present, there has not been a nusenipt No. 188266. Received September 14. 2000; approv suitable method to measure the fracture toughness of laminated This work was supported by the National Science Foundation of China ceramics because these methods are suitable only for homoge- neous ceramics while laminated ceramics are not homogeneous i
journal J Am. Cetiiin. S-i . HS |l(l| 24'i7-61 (2002) Control of Composition and Structure in Laminated Silicon Nitride/Boron Nitride Composites Chang-ati Wang,* Yong Huang,* Qingfeng Zan. Linhua Zou, and Shengyou Cai State Key Labonilory of New Ceramics and Fine Processing. Department ol Materials Science ;ind Engineering. Tsinghua University. Beijing. I(X)O84. People's Republic nl" China Based on a biomimetic deNign, Si^N^/BN compasites with luminated structures have been prepared and investigated through i-i)mpo.siti(>n control and structure design. To further improve the mechanical properties of the compo.sites. Si,N4 matrix layers were reinforeed by SiC whiskers and BN separating layers were modified hy adding Si^Nj or AUO,. The result.s showed that the addition of SiC whiskers in the SiiN^ matrix layers eould greatly improve the apparent fracture toughness (reaching 28.1 MPa ni"'). at the same time keeping the higher hendiug strength (reaching 651.5 MPal of the composites. Additions of 50 wt% .XUO^ or 10 wt% Si^Nj to BN interfacial layers had a beneficial effecl on the strength and toughness of the laminated Si^Nj/BN composites. Through observation of microslructure by SEM. multilevel toughening mechanisms contributing lo high toughness of the laminated Si^Nj/BN composites were present as the first-level toughening mechanisms from BN interfacial layers as crack deflection, bifurcation, and pull-out of matrix sheets, and tbe secondary lougbening meehanism Irom whiskers in matrix layers. I. introduction S ILICON NiTRiDH is known as an important high-temperature structural ceramic with excellent mechanical propenies from anibienl lo elevated temperatures, but the brittleness limit.s its applications. Therefore, various approaches have been cbosen to increase its toughness including self-reinforcement' and incorporation of particles and whiskers,'""* etc., but improvement in the toughness of the material has been only limited. However, it is well known that some natural biomaterials, such as bamboo, wood, and nacre, have specific composite structures and exhibit good mechanical properties in that they have high toughness and excellent damage resistance, compared with artificial composites.^ A typical example of a layered biomaterial is mollusk shell. The hierarchical structure of a mollusk shell is formed by aragonite layers about I fxm thick joined by a kind of mortar of proteins. This particular configuration imparts over I order of magnitude higher bending strength and toughness than those of aragonite single crystals. These natural composites may give us some insights into making better structural materials with high toughness through biomimetic design. Recently, some researchers introduced laminated or multilayered structures to improve the mechanical properties of stmctural ceramics. In 1990. Clegg'* reported preparing a laminated SiC compttsite separated by graphite layers, and the apparent fracture toughness of the composite was improved from 3.6 to 17.7 N. Clausseii—conlribuling editor Manuscripl No. I8K266. Received Sepleniber !4. KHXk approved February 12. m. This work was suppined by ihc National Science Foundution of China. 'Member. American Ceramic Society. MPa-m''". Thereafter, various methods including tape casting and slip casting have been applied to make ceramic-matrix composites with multilayered structtircs. The laminated design is considered to be a promising method to overcome the brittle nature of ceramics. However, the bending strength of the laminated materials at mom temperature is considerably reduced because ot the presence of weak separating layers.^'"'^ Laminated composites consist of ceramic-matrix and interfacial separating layers. The former mainly determines the bulk strength and the latter gives rise to the high toughness due to the deflection of crack. Therefore, to improve ihe mechanical properties of these composites, two ways may be used: adjusting the composition and structure of the inierfacial separating layers and/or reinforcing Ihc ceramic matrix layers. In the current paper, laminated SiiNj/BN composites were prepared to imitate the structure of mollusk shell: thin Si,N4 layers were laminated like the aragonite plates with BN separating layers as the organic joining. To improve ihe mechanial properiies of laminated Si^fN^BN composites. SiC whiskers were added to Si ,N^ matrix layers for further reinforcement of matrix layers, and Ai,O, or Si,N4 was Incorptiraied into the BN layers for the purpose of adjusting the bonding strength between S'y^N^ layers. II. Experimental Procedure SiiN4 {Founder High-Tech Ceramic Corp.. China) powders with 7 wi% Y,0, (99.9%. Hokke Chemicals. Japan) and 3 wt% AI,O, (purity > 99.9<7f) were ball-milled with 20 wi9i SiC whiskers (TWS-4()(). Tokai Carbon Co.. Japan) in ethanol for 24 h to achieve a homogeneous mixture. The slurry was dried and sieved through a 60-mesh screen; then the powder was mixed with organic binders and then rolled into a sheet about 200 fini thick by rolling compaction. The green sheets were subsequently coated with a slurry of BN adjusted with AUO, or SiiNj. and then dried and packed into a graphite die. After the binders were burned out at 430"C (bi- 2 h in air, the green body was hot-pressed in a graphite resistance furnace in flowing N, al I82O''C for \.5 h. For comparison, SiiNj monolithic samples were fabricated using the same composition and processing of Si,N_,/BN laminated composites, but without the BN coating on ihe Si,Nj green sheets. The mechanical properties of the laminated Si^N/BN composites were determined on an A-2000 Shimad/u universal materials testing machine. Six samples were tested for each property. The loads were applied perpendicular to the plane of laminates. Bending strength was determined by three-point bend testing (test bars 4 mm x 3 mm X Mt mm) with a span of 30 mm and a crosshead speed of 0.5 mm/min. The tensile surface of the samples was polished with diamond paste down to I |xm. and the corners were chamfered with a l^i mm diamond grinding wheel. As for the fracture toughness of laminated ceramics, even though there are only a very few references mentioned, in which the fracture toughness was measured by SENB.'' SEVNB,'" and the indentation method.'* up to the present, there has not been a suitable method to measure the fracture toughness of laminated ceramics because these methods are suitable only for homogeneous ceramics while laminated ceramics are not homogeneous in 2457
Journal of the American Ceramic Sociery-Wang et al. Vol. 85. No microstructure. For comparison. the same SENB method was used to measure the fracture toughness of monolithic and laminated ceramics in this paper. For the reasons mentioned above, we used apparent fracture toughness rather than fracture toughness for laminated ceramics. The dimension of test bars was 4 mm x 6 mm X 30 mm. The testing was conducted with a crosshead speed of 0.05 mm/min and a span of 24 mm. The load-displacement urve was also determined for notched samples using the same conditions as those used in the determination of fracture toughness To assess the bonding strength of BN interfacial layers, a so-called double-shearing method was designed as shown in Fig 1. The samples were made with three layers of thick Si, N, layers and two BN layers using the same compositions and processing of laminated Si, N,/BN composites The microstructure of the laminated Si, N /BN composites was observed by scanning electron microscopy (SEM). The surface of a Si, N, matrix layer was polished to a mirror surface and etched with melted NaOH Ill. Results and Discussion 250um (1) Microstructure and Load-Displacement Curve of Laminated Si,N,/BN Composites Figures 2(a) and (b) show SEM photographs of the surfaces perpendicular and parallel to the stacking sheets, respectively. In Fig. 2(a), the laminated structure can be observed clearly, in which SiNa layers(black) are about 80-100 um thick and BN(white) separating layers are only-20 um thick. Figure 2(b) shows that in the Si,N4 matrix layers SiC whiskers are distributed homoge neously and preferentially oriented due to rolling compaction an hot-pressing. In the composites, the thickness of the matrix Si,N thickness of the interfacial BN layers depends on the viscosity of the BN slurry and coating times. Therefore, the thickness ratio between matrix SiaNa layers and interfacial BN layers can be easily controlled and adjusted The above specific microstructure makes this kind of laminated Si,N,/BN composite with the specific mechanical performance Figure 3 shows two typical load-displacement curves of laminated Si N/BN compo site and conventional monolithic Si, N, material respectively. It can be seen that the laminated Si, N,/BN composite exhibits a nonbrittle failure manner while the conventional mono- lithic Si,N, material fractures catastrophically. For a laminated 9 H Si,N,/BN composite, after the first load drop, the load-bearing ability of the tested laminated ceramic bar still retains over 50% of the peak load. Until totally fractured. this laminated Si N /BN Fig. 2. SEM photographs of the surfaces of laminated Si,N /BN com- composite gives a prolonged deflection besides the elastic defor- posites:(a) side view perpendicular to the stacking sheets and (b) matrix mation. This shows that a laminated composite exhibits a different Si, N, layer, in which whiskers were aligned parallel to the stacking sheets. fracture behavior from that of a monolithic ceramic. Hence it may Load laminated composite Sin Fig. I. Schematic chart of double-shearing method for determinat of bonding strength of BN interfacial layer in laminated Si,N 图N Fig 3. Typical load-deflection curve of laminated Si, N, /BN composites
2458 Journal of the American Ceramic Society—Wang et al. Vol. 85. No. 10 microstruclure. For comparison, the same SENB method was used to measure the fraciure toughness of monolithic and laminated Leramics in this paper. For the reasons mentioned above, we u.sed apparent fracture loughness rather than fracture toughness for laminated ceramics. The dimension of test hars was 4 mm X 6 nmi X 30 mm. The testing was conducted with a crosshead .speed of 0.05 mm/min and a span of 24 mm. The load-displacement curve was also determined for notched samples using the same conditions as those used in the detemiination of fracture toughness. To assess the bonding strength of BN interfacial layers, a so-called double-shearing method was designed as shown in Fig, I, The samples were made with three layers of thick Si ,N^ layers and two BN layers using the same compositions and processing of laminated Si^Nj/BN composites. The microstructure of the laminated Si^Nj/BN composites was observed by scanning electron microscopy (SEM). The surface of a Si,N.j matrix layer was polished to a mirror surface and etched with melled NaOH, in. Results and Discussion (/) Microstructure and iMad-Displacement Curve of laminated Si^NyBN Composites Figures 2(a) and (b) show SEM photographs of the surfaces perpendicular and parallel to the stacking sheets, respectively. In I-ig. 2{a). the laminated structure can be observed clearly, in which Si,Nj layers (black) are about 80-100 |im thick and BN (white) separating layers are only -20 p.m thick. Figure 2(b) shows that in the Si,Nj matrix layers SiC whiskers are distributed homogeneously and preferentially oriented due to rolling compaction and hot-pressing. In the composites, the thickness of the matrix Si,N4 layers depends on the thickness of green SiiN,j sheets, while the thickness of the interfacial BN layers depends on the viscosity of the BN slurry and coating times. Therefore, the thickness ratio between matrix Si^tNj layers and interfaeial BN layers can be easily controlled and adjusted. The above specific microstruclure makes this kind of laminated .Si,N4/BN composite with the specific mechanical performance. Figure 3 shows two typical load-displacement curves of laminated Si.N^/BN composite and conventional monolilhic Si^Nj material. respectively. It can be seen that the laminated SiiNj/BN composite exhibits a nonbrittle failure manner while the conventional monolithic Si,N., material fractures catastrophically. For a laminated SijNj/BN composite, after the first load drop, the load-bearing ability of the tested laminated ceramic bar still retains over 50% of the peak load. Until totally fractured, this laminated Si^Nj/BN composite gives a prolonged deflection besides the elastic deformation. This shows that a laminated composite exhibits a different fracture behavior from that of a monolithic ceramic. Hence it mav Load • 566666^ V V V V V V S BN Support Fig. 2. SEM photographs of the surfaces of laminated Si^Nj/BN composites: (a) side view perpendicular to the stacking sheets and (h) matrix Si,Nj layer, in which whiskers were aligned parallel to ihc stacking sheets. 600 400 laminated composite 100 200 300 Displacement V\^. I. Schematic charl of double-shearing method for determination nf honding strength of BN interfacial layer in laminated Si3N4/BN composites. 400 Fig. 3. Typical load-defleetion curve of laminated SijNi/BN composites with notch
October 2002 Control of Composition and Structure in Laminated Si,N/BN Composites 2459 Table L. Mechanical Properties of L aminated Si, N,/BN Composites with Different Reinforcements in Si,N, Matrix Layers Secondary Apparent fracture toughness Bending strength MPam (MPa Remarks 7.3±0.5 2.2* 34.5 Monolithic Si Na ceramics SiC whiskers 8.4±0.8 0.7+ 99.7 Monolithic SiC(w ySi, N, composites (20w%) 75 wt%O BN 5.1±1.1 498.4 *22.7 Laminated SiaN,/BN composites 25 wt%AL-O, SiC whiskers 75 wt%o BN 28.1±1.0 651.5+ 74.9 Laminated Si,N,/BN composites +25 Wt%AL,O be inferred that the mechanical performance of ceramic materials The conbination of the two-level toughening mechanisms contrib- can be substantially improved by a special structural design similar utes to the high toughness of the laminated Si N,/BN composites to that conducted on biomaterials, as already indicated in the Another interesting result should be noted that SiC literature, 0, II increase both the bending strength and the apparent fracture toughness of the laminated Si, N,/BN composites to a greater (2) Reinforcement of Ceramic Matrix Layers by extent than it does in the monolithic Si,N, ceramics; i.e., the SiC Whiskers influence of whiskers is higher for Si,N, materials with a multi The mechanical properties of Si, N,/BN composites with whis- layered structure than for monoliths. The reason for that has not er reinforcements added in Si, N, matrix layers are summarized in been very clear up to now. The effects of SiC whiskers may be Table 1. The composition of the separating layer was 75 wt% BN attributed to the great enhancement of resistance to crack initiation wt%e AlO and propagation. Generally, by comparing samples I and 3(or 2 and 4), it is obvious that the introduction of weak interfaces(BN Al,O (3) Adjusting Composition of Interfacial Layers nto the Si,N, matrix can improve the fracture toughness As discussed above, the laminated structure mainly contribute from 7.3 to 15.1 MPam"(or from 8.4 to 28. 1 MPam). This to high toughness of the laminated Si N,/BN composites. The shows that the laminated structure with weak interfacial layers is main mechanism is that the interfacial layer deflects the crack he main reason for the high toughness of laminated SiN/BN epeatedly and consumes a large amount of fracture energy when composites. However, improvement in the apparent fracture he material is loaded. To get the maximum toughening effect, the toughness is usually at the expense of the decreased bendi bonding of interfacial layers must be moderate and suitable. If the strength of the composites because a large number of defects are interfacial layer is too weak, it will dissipate less energy, although Sic whiskers into the Si,, matrix layers ly enhances the Furthermore, it may lower the strength of the laminated Si,N,/BN apparent fracture toughness of the laminated Si, N /BN composites composites. On the other hand, if the interfacial layer is too strong. from 15. I to 28. 1 MPa m while at the same time, the bending he composite will behave as a brittle monolithic ceramic and may strength is retained at 651 MPa, greater than that of the laminated have a low apparent fracture toughnes composites without SiC whiskers (498 MPa). Therefore, the Since BN powder is unsinterable at the sintering temperature of conclusion can be drawn that weak interfacial layers increase the Si, Na, the interfacial layer with pure BN forms a very weak apparent fracture toughness but decrease the bending strength while SiC whiskers as reinforcements in the Si,N, matrix lavers can preserve the bending strength and increase the apparen erfacial layers as separating layer modifiers. fracture toughness simultaneously at a relatively high level. The A) Adding Al, O, as a Separating Laver Modifier: Table II ceramic matrix layers determine the total strength of laminated and Fig. 5 show the effects of Al O, as a separating layer modifier omposites with an interfacial separating layer with a fixed on the bulk mechanical properties of laminated Si,N/BN com- composition: i.e., the stronger the matrix layer, the higher the sites. The results show that with the increase of Al,O,content strength of the laminated composites. n BN layers the bending strength and apparent fracture toughness Figure 4 shows crack propagating paths in the laminated of the composites increase to a maximum and then begin to Si, N/BN composites. The crack from SiaN, matrix layers is decline. At a composition of 50 wt% BN 50 wt% Al,O, for deflected into the BN separating layers and propagates along the interfacial layers, both the bending strength and the apparent BN layer repeatedly as shown in Fig. 4(a). Besides the deflection fracture toughness of the composites reach the maximum values of crack from BN layers. there are crack bifurcation( Fig. 4(b))and Therefore, it can be concluded that the composition of 50 wt% BN matrix layer pull-out(Fig. 4(c)) toughening mechanisms in the 50 wt% Al,O, for interfacial layers may be optimal for Al-O, composites. These toughening mechanisms from the interfacial layers are called the first-level toughening mechanisms and mainly During sintering of the laminated Si, N, /BN composites, the contribute to high toughness of the laminated composites sintering aids in matrix Si,Na layers including Y,O,, Al,O3, and It is well known that whiskers can reinforce ceramics to some esidual SiO, on the surface of Si,N, powder will diffuse into the extent by means of crack deflection, whisker bridging and pulling BN interfacial layer. The diffused additives will react with Al,O out, and other reinforcing mechanisms. Especially in the case of in the interfacial layer and form a eutectic glassy phase which whisker preferred orientation in ceramic matrix, the whisker ncreases the bonding of the interfacial layer. At a lower amount of reinforcement effect becomes more remarkable, 27.18 As men Al,O, doped in the BN layer, the amount of the glassy phase is too tioned above, SiC whiskers are two-dimensionally distributed in little and leads to a weaker bonding interface; therefore, the Si, N, matrix layers, so the effect of whisker toughening is more strength and toughness of the composites are lower. With an markable; i.e., whiskers in matrix layers further toughen the ease in the amount of Al, O, doped in the BN interfacial lay smpes ites, bect os i ter ciane aovt w hihe w hiskey s hening is the hindang oo thnt ot glassy ayase howeyestrwer the al. o alled as the secondary-level toughening mechanism. content in the BN interfacial layer is so high that Al,O, becomes The above mechanical property results suggest that the dominant component of the interfacial layer, the interface exist a synergic reinforcing effect between the whisker becomes Al, O, ceramic. Since there is a large difference between ment in the matrix layers and weak interfacial layer tou the thermal expansion coefficients of Al,O,(8.9 x 10/K
October 2002 Control of Composition and Structure in Laminated SI,NyBN Composites 24,59 Table I. Mechanical Properties of Laminaled Si3N4/BN Composites with Different Reinforcements in SifNj Matrix Layers Sec I mil a ry Sainpk' reinlbrLement Separating layer composition Appiiient Irai'ture tdug M P " (MPa) Retnarks 1 T 3 4 — SiC whiskers (20 wt%) — SiC whiskers (20 Wt^r) 75 wt% + 25 75 wl% + 25 — — BN wt'^AUO, BN wt%AUO, 7,3 8.4 15.1 28.1 ±0.5 ±0.8 ± 1.1 ± 1,0 862,2 ± 34.5 820.7 + 99.7 498.4 ± 22.7 651.5 i 74.9 Monolithic Si,N., ceramics Monolithic SiC(w)/Si,N4 composites Laminated Si3N4/BN composites Laminated Si,N,/BN composites be inferred that the mechanical performance of ceramic materials can he substantially itiiproved by a special structural design similar to ihat conducted on biomaterials, as already indicated in the literalure.'"-" (2) Reinforcement of Ceramic Matrix Layers by SiC Whiskers The mechanical properties of Si,N4/BN composites with whisker reinforcements added in SijNj matrix layers are summarized in Table 1. The comptisition of the separating layer was 75 wt% BN + 25 wt% AUO;,, Generally, by comparing samples I and 3 (or 2 and 4), it is obvious thai the introduction of weak interfaces (BN + AUO,) into the Si,N4 matrix can improve the apparent fracture toughness from 7.3 lo 15.1 MPa-m"- (or fr<)m 8.4 to 28,1 MPa-m""), This shows that ihc laminated siiitcture with weak interfacial layers is the main reason for the high toughness of laminated Si^N/BN cotiiposites. However, improvement in the apparent fracture toughness is usually at the expense of the decreased bending strength of Ihe composites because a large number of delects ai"e introduced into the composites by the weak interfacial layers. Comparison of samples 3 and 4 shows that the incorporation of SiC whiskers into the Si,N4 matrix layers greatly enhances the apparent fracture toughness of the laminated Si,N4/BN cttmposites from 15,1 lo 28.1 MPa-m""^. while at the same time, the bending strength is retained at 651 MPa, greater than thai of the laminated composites without SiC whiskers (498 MPa), Therefore, the conclusion can be drawn that weak intertVcial layers increase the apparent fracture toughness but decrease the bending strength, while SiC whiskers as reinforcements in the Si^Nj matrix layers can preserve the betiding strength and increase the apparent fracture toughness simultaneously at a relatively high level. The ceramic matrix layers determine the itital strength of laminated composites wilh an interl'acial separating layer with a fixed CO til position; i.e.. the stronger the matrix layer, the higher the strength of the laminated composites. Figure 4 shows crack propagating paths in the laminated Si,N4/BN composites. The crack from Si,N4 matrix layers is deflected into the BN separating layers and propagates along the BN layer repeatedly as shown in Fig. 4(a). Besides the deflection of crack from BN layers, there are crack bifurcation (Fig. 4(b)) and matrix layer pull-oul (Fig. 4(c)t toughening mechanisms in the composites. These toughening mechanisms from the interfacial layers are called the first-level toughening mechanisms and mainly contribute to high toughness of the laminated composites. It is well known that whiskers can reinforce ceramics to some extent by means of crack deflaction. whisker bridging and pullingout. and other reinforcing mechanisms. Especially in the case of whisker preferred orientation in ceramic matrix, ihe whisker reinforcement effect becomes more remarkable."^''•"* As mentioned above. SiC whiskers are two-dimensionally distributed in SiiNj matrix layers, so the effect of whisker toughening is more remarkable: i.e.. whiskers in matrix layers further toughen the composites. Because the zone of whisker acting is relatively smaller than thai of interfacial layers, the whisker toughening is called as the secondary-level toughening mechanism. The above mechanical property results suggest that there may exist a synergic reinforcing effect between the whisker reinforcement in the matrix layers and weak intertacial layer toughening. The conbination of the two-level toughening mechanisms contributes to the high toughness of the laminated SiiN4/BN composites. Another interesting result should be noted that SiC whiskers increase bolh the bending strenglh and the apparent fracture toughness of the laminated Si,N4/BN composites to a greater extent than it does in the monolithic Si,Nj ceramics; i.e.. the influence of whiskers is higher for Si,N4 materials with a multilayered structure than for monoliths. The reason for that has not been very clear up to now. The effects of SiC whiskers may be attributed to the great enhancement of resistance to crack initiation and propagation. (3) Adjusting Composition of Interfacial iMyers As discussed above, the laminated structure mainly contributes to high toughness of the laminaled Si,N4/BN composites. The main mechanism is that the interfacial layer deliects the crack repeatedly and consumes a large amount of fracture energy when the material is loaded. To get the maximum toughening effecl. the bonding of interfacial layers must be moderate and suitable. If the interfacial layer is too weak, it will dissipate less energy, although cracks can easily be deflected along the interfacial layer, and this case is unf'avorable to the toughness of the laminated composites. Furthertiiore, it may lower the strength of ihe latiiinated Si,N4/BN composites. On the other haiul, if the interfacial layer Is too strong, the composite will behave as a brittle monolithic ceramic and may have a low apparent fracture toughness. Since BN powder is unsinterable at the sintering temperature of Si3N4. the intertacial layer with pure BN fornis a very weak interface between Si3N4 matrix layers. To adjust the bonding strength of the interfacial layers, AUO, or Si ,N4 was added to BN interfacial layers as separating layer modifiers. (A) Adding AI2O, as a Separating Layer Modifier: Table II and Fig, 5 show the effects of AUO, as a separating layer modifier on ihe bulk mechanical properties of laminated SitN^/BN composites. The results show that wilh the increase of AUO, content in BN layers the bending strength and apparent fracture toughness of the composites increase to a maximum and then begin lo decline. At a composition of 50 wl9c BN 4- 50 wf^ AUO, for tnterfacial layers, both the bending strength and the apparent fracture toughness of the composites reach the maximum values. Therefore, it can be concluded ihat the composition of 50 wt% BN + 50 wt% AUO, for interfacial layers may be optimal lor AUO, as a separating layer modifier. During sintering of the latninated Si,N4/BN eomposite.s. the sintering aids in matrix Si,N4 layers including Y,0,. AUO,. and residual SiO^ on the surface of Si,N4 powder will diffuse into the BN interfacial layer,''^ The diffused additives will react with AUO-, in the interfacial layer and form a eutectic glassy phase which increases ihe bonding of the inlerfacial layer. Al a lower amount of Al^O, doped in the BN layer, the amount of the glassy phase is too little and leads lo a weaker bonding interlace; therefore, ihe strength atid toughness of the composites are lower. With an increase in the amtiunt of AUO, doped in the BN interfacial layer. the bonding of the interfacial layer becomes stronger because of the higher amount of glassy phase. Htiwever. when the AUO3 content in the BN interfacial layer is .so high that AUO, becomes the dominant component of the interfacial layer, the interface becomes AUO, ceramic. Since there is a large difference between the thermal expansion coefficients of AI^O, (8.9 X 10 ^/K
Jounal of the American Ceramic Society-Wang et al. (a) Table II. Mechanical Properties of the Laminated Si, N,/BN Composites with Al,O, as Separating Layer Modifier ALO 18.7±2.04659±43.3 24.9±1.6623.1±27 28.9±4.17095±89 23.4±2.0651.5±74.9 AL,O 100 7.7±1.8666.8±37.9 toughness of the composites are the highest. The above expe ments show that this optimal interface composition is 50 wt% BN +50 wt% AlO DO um (B) Adding SiNN, as a Separating Layer Modifier: Table Ill modifier on the bulk mechanical properties of laminated Si, N,/BN (b) composites. The results show that Si,N, as a separating layer modifier appears a little bit different from Al,O,. With an increase of Si N, content in the BN layers, the bending strength of the composites increases and reaches a maximum at 75 wt% Si, Nas while the apparent fracture toughness of the composites reaches a maximum at 10 wt% SiN, and then decreases The effect of Si Na doping in the BN interfacial layer is different from that of Al,O, doping. During sintering of the laminated Si N,/BN composites. the sintering aids in the Si, N matrix layer diffuse into the BN interfacial layer and assist with sintering the Si N4 doped in the BN interfacial layer. With an increase of Si Na content in the BN interfacial layer, the sintering improves and thereby the bonding strength of the interfacial layer increases gradually. When the content of Si,N, in the BN interfacial layer is close to 100 wt%, the interfacial layer is almost the same as the matrix layer, and the laminated composite behaves as a monolithic brittle Si,Na ceramic. Figure 7 shows the effect of Si,N, content in the BN interfacial layer on the shear strength of 80 um he interfacial layer. It can be found that the shear strength of the interfacial layer linearly increases with respect to the content of Si,Na in the BN interfacial layer. In the case of this interfacial state, the stronger the interfacial layer, the higher the strength of the composites, but at the expense of the apparent fracture toughness of the composite. As a compromise, about 10 wt% E600 180pm Fig. 4. SEM photographs of crack propagating paths in laminated SiN/BN composites showing (a) crack deflection. (b) crack bifurcate and (e)matrix layer pull-out. 20-1000oC)and Si, N, (3.2 X 10"/K 20-1000 C), the interfa- cial layer is under tensile stress after cooling from the sintering AL,O, content in BN separating layer(wt%) state, which is unfavorable for strength and toughness of the composite. Therefore, there exists an optimal moderate amount of Al O, doped in the BN interfacial layer, at which the strength and with AL, O, as a separating layer modifierhinated Si,N,/BN Fig. 5. Mechanical properties of the lar
2460 Journal of the American Ceramic Societ}-—Wang et al. Vol. 85, No. 10 Table II. Mechanical Pniperties of the Laminated Si^N/BN Composites with AI,O, as Separatin};; Layer Modifier Fig. 4. SEM photographs of crack propagiiiing paths in laminated SiiNj/BN coiiiposites showing (a) crack detlectioii, (b) crack biturcation, and (c) matrix layer pull-out. 20-I000"C) and Si.N^ (3,2 X lO'^/K. 20-1000°C). the interfacial layer is under ten.sile stress after cooling from the sintering state, which is unfavorable for strength and toughness of the coniposite, Theretore. there exists an optimal moderate amount of AliO, doped in the BN intertacial layer, at which the strength and Sample Dopant in BN sepiirrttjng iiiyer Dopani conieni in BN Mjparaiing luver (wt%) Apparcni fraciure toughness (MPa-m"'') Bending strength (MPa) AI2O3 0 25 50 75 100 18,7 24.9 28,9 23.4 7,7 ± 2,0 ± 1.6 ±4.1 ±2.0 ± 1,8 465.9 ± 43. 623.1 ± 27, 709.5 ± 89, 651.5 ± 74. 666.8 ± 37, 3 1 6 9 9 toughness of the composites are the highest. The above experiments show that this optimal interface composition is 50 wt% BN + 50 wt% AUO3. (B) Adding Si^N^ as a Separating Layer Modifier: Table III and Fig. 6 show the effects of Si,N4 as another separating layer modifier on the bulk mechanical properties of laminated Si ,N^/BN composites. The results show that Si,N4 as a separating layer modifier appears a little bit different from AUO^. With an increase of Si-,N4 content in the BN layers, rhe bending strength of the composites increases and reaches a maximum at 75 wtVr Si,N^. while the apparent fracture toughness of the composites reaches a maximum at 10 wt'Jf SiiN4 and then decreases. The effect of Si,N4 doping in the BN interfacial layer is different from that of AUO^ doping. During sintering of the laminated Si,N4/BN composites, the sintering aids in the SiiN^ matrix layer diffuse into the BN interfacial layer and assist with sintering the Si^Nj doped in ihe BN interfaciat layer. With an increase of Si^Nj content in the BN interfacial layer, the sintering improves and thereby the bonding strength of the interfacial layer increases gradually. When the content of SiiNj in the BN interfacial layer is close to 100 wf^/r. the interfacial layer is almost the same as the matrix layer, and the laminated composite behaves as a monolithic brittle .Si^Nj ceramic. Figure 7 shows the effect of SiiNj content in the BN interfacial layer on the shear strength of the interfacial layer,"" It can be found that the shear strength of the interfaeial layer linearly increa.ses with respect to the content of Si3N4 in the BN interfacial layer. In the case of this interfacial state, the stronger the interfacial layer, the higher the strength of the composites, but at the expense of the apparent fracture toughness of the composite. As a compromise, about 10 wt% OJD 2 •5 c 900 800 700 600 500 CO 400 - Hu ^^ es b y 0 20 40 60 80 100 Al^Oj content in BN separating layer (wt%) Fig. 5. Mechanical properties of the laminated Si,N4/BN composites with AUO, as a separaUng layer mtKlifier
October 2002 Control of Composition and Structure in Laminated Si,N/BN Composites Table lll. Mechanical Properties of the Laminated Si,N,/BN Composites with Si, N4 as Separating Layer Modifier 18.7±2.0465.9±43.3 00165840739 钥20 80±0.9649.7±38.7 7.8±0.61124.6±143.2 7.1±0.3970.0±13.5 102030405060708090100 doped in the BN interfacial layer may be suitable for anical properties of the laminated Si,N,/BN composites, as raction of si, N, in BN interfacial layer in the above results Fig. 7. Shear strength of BN interfacial layer versus the fraction of Si,Na doped in the BN interfacial layer. IV. Summary and Conclusions Based on a biomimetic design, Si, N, /BN composites with laminated structure can be easily prepared using a simple tech- whisker toughening in matrix layers may exist as a secondary nique, rolling compaction and hot-pressing. This laminated struc toughening mechanism, which also contributes to the high tough ture results in very high toughness of the composites in the ness and good strength of the composites orientation tested Because of the presence of weak interfacial layers, the bending References trength of the Si N /BN composites often decreases despite high toughness. The addition of SiC whiskers in Si,N, matrix layers car N. Hirosaki. Y. Akimune, and M. Mitomo, " Effect of Grain Growth of B-Silicon greatly improve both the bending strength (651.5 MPa) and the Nitride on Strength, Weibull Modulus, and Fracture Toughness. "J Am. Ceran. Soc. toughness (28. 1 MPam)of laminated Si N/BN composite 76711892-94(1993) Al,O, and Si, Na are two effective modifiers for interfacial SiC-Whisker- Reinforced S. bricatod by Extrusion画HPe图Am layers of laminated Si, N,/BN composites. Adding 50 wt% Al,O Ceram Soc,76|611420-24(1993 or 10 wt% Si, N, to BN interfacial layers makes interfacial layers S. T. Buljan, J. G Baldoni, and M. L. Huckabee. "Si, N,-SiC Composites."A with moderate bonding strength and results in laminated Si3N,/BN eran Sot.Bal.6612347-52(1987 G. Pezzotti,"Effect of HIP Sintering on the Crystal Structure and Fracture composites with good strength and toughness aviour of a-SiC Platelets Embedded in Si, Na Matrix. "J Ceran. Soc. Jpn. 101 There exist multilevel toughening mechanisms contributing to 7811131-36(199 high toughness of laminated SiaN,/BN composites, Crack deflec 'B. L Zhou, "The Biomimetic Design of Worst Bonding Interface for Ceramic tion, bifurcation, and pull-out of matrix sheets are suggested as Matrix Composites, Compos. Eng, 5 [10-11 1261-73(1995) w. J. Clegg. K Kendall, N. M. Alford, D, Birchall, and T. w. Buton,"A Simple first-level toughening mechanisms due to the modified BN inter- Way to Make Tough Ceramics, " Nature(London), 347 [101 445-47(1990 facial layers and the main reasons for high toughness, while L Zhang and V D. Krstic, "High To raphite Laminar Composite by Slip Casting, Theor, Appl. Mech,24.13-19(1995) J, Phillipps. w. J. Clegg, and T. w. Clyne."Fracture Behavior of Ce 1400 M4817xmM等 Experimental Data. "Acta red Silicon Nitride, J, An, Ceram, Soe. 1000 IT. Chartier, D, Merle, and J. L. Besson, "Laminar Ceramic Composites, "J 曾60xx I Y Huang H N. Hao, Y. L Chen, and B N. Zhou, "Design and Preparation of Silicon Nitride Composite with High Fracture Toughness and Nacre Structure, Acta all Sin,9161479-8401996) " H. Liu and S. M. Hsu, ""Fracture Behavior of Multilayer Silicon Nitnde/Boron 200 Nitride Ceramics, "J. A Ceran. Soc., 79 1912452-57(1996) E. H. Lutz and M. V. Swain, "Fracture Toughness and Thermal Shock Behav of Silicon Nitride- Boron Nitride Ce " J. A Ceram Soc,7511167-70(1992) 28 AD, Kovar, M. D. Thouless, and J, W. Halloran, "Crack Deflection and Propaga- in Layered Silicon Nitride/Boron Nitride Ceramics, "J, Am. Ceram. Soc., 8 20 Z Lences, K Hirao, M. E. Brito, M. Toriyama, and S, Kanzaki, "Layered Silice Nitride-Based Con ith Discontinuous Boron Nitride Interlayers,J 12 eT, Ohji, Y. Shigegaki, N, Kondo, and Y. Suzuki, "Fracture Toughness of YE Multilayer Silicon Nitride with Crack Deflection, "Mater, Lert, 40, 280-84(1999 C A. Wang, Y. Huang, and Z P. Xie, "The Effect of Whisker Orientation in SiC whisker-Reinforced Si, N4 Ceramic Matrix Composites."J. Eur. Ceram. Soc.. 19[101 199(99 Silicon Carbide-Whisker-Reinforced Silicon Nitride- Matrix Composites by Whisker- Oriented Alignment. "J Am. Ceram. Soc., 84 [I 161-64(2001). Si, N, content in BN separating layer (wt%) Fig. 6. Mechanical properties of the laminated Si,N,/BN composites of Laminated Si N4 Matrix Ceramics": Ph. D, Dissertation. Tsinghua Universit with Si, N, as a separating layer modifier Beijing, People's Republic of China, 1998
October 2002 Control of Composition and Struclure in Laminated Composite.^ 2461 Table III. Mechanical Properties of the Laminated Compnsite.s with Si,N4 as Separating Layer Modifier Sample 1 3 4 5 6 Dopanl in BN separating liiycr Si,N, SijN^ Si3N4 Si,N^ Si,N4 Dopiini ctinteni in BN scparaiiiig layer (wi%) 0 10 25 50 75 100 Apparenl frdtiure toughness (MPa-m"') 18.7 ± 2.0 20.2 ± 2.6 7.9 ± i.6 8.0 ± 0.9 7.8 ± 0.6 7.1 ± 0.3 Bending strength (MPa) 465.9 ± 43.3 644.5 ± 56.2 554.0 ± 73.9 649.7 ± 38.7 1124.6 ± 143.2 970.0 ± 13.5 SiiNj doped in ihe BN intertacial layer mechanical properties of thie iaminated Si^ shown in the above resulls. may be suitable for coniposites. as IV. Summary and Conclusions Based on a biomiinetic design. Si^Nj/BN composites with laminated structure can be easily prepared using a simple technique, rolling compaction and hot-pressing. This laminated structure results in very high toughness of the composites in the orientation tested. Because of the presence of weak interfacial layers, the bending strength of the Si,N4/BN composites often decreases despite high loughness. The addition of SiC whiskers in Si ,N4 matrix layers can greatly improve both the bending strength (651.5 MPa) and the toughness (28.1 MPa-m''") of laminated Si,N4/BN composites. AUO, and Si,N4 are twt) effective modifiers for interfacial layers of laminated Si^N^/BN composites. Adding 50 wt% AUO;, or 10 wt% Si3N4 to BN inlerfacial layers makes interfacial layers with moderate bonding strength and results in laminated Si,N4/BN composites with good strength and toughness. There exist multilevel toughening mechanisms contributing to high toughness of laminated SitNj/BN composites. Crack deflection, bifurcation, and pull-out of matrix sheets are suggested as first-level toughening mechanisms due to the modified BN interfacial layers and the main reasons for high toughness, while 0 20 40 60 80 100 Si N^ content in BN separating layer (wt%) Fig. 6. Mechanical properties of the laminated SijN4/BN composites with Si3N4 as a separating layer tnodifier, • « 40 • S 20 - 0 10 20 30 40 50 60 70 80 90 100 Fraction of Si^N^ in BN interfacial layer Fig. 7. Shear strength of BN interfacial layer versus the fraetion of Si.iN^ doped in fhe BN interfacial layer. whisker toughening in matrix layers may exist as a secondary toughening mechanism, which also contributes to the high toughness and good strength of the composites. References ' N. Hiro.saki. Y. Akimune, und M. Mitomo, "Effeet of Grain Growth of p-Silicon Nitride on Strength. Weibull Modultjs. and Fracture Toughness," J. Am. Ceram. Soc., 76|7 | 1892-94 11993). -Y. Goto and A. Tsuge, "Mechnical Properties of Unidirectional I y Oriented SiC-Whisker- Reinforted Si^Nj Fabricated by Extrusion and Hol-Prcssing." /. Am. Ceram. Soc.. 76 [6| 1420-24 (1993). 'S. T. Buljan. J, G. Baldoni. and M. L. Huckabee. "SijNi-SiC Composites." Am. Ceram. Soc. Bull.. 66 [2| 347-.*i2 (1987), •"G, Pe/./otii, "Effect of HIP Sintering on the Crystal Sinicturc and Fracture Behiiviour of a-SiC Pliitelets Embedded in Si,Nj Matrix," J. Ceram Soc. Jpii.. IDI |I178| 1131-36(1993). 'B. L. Zhou. "The Biomimeiic Design of Worst Bonding Interface for Ceramic Mulrix Comptisites.'" Com/Kw. En^t.. S \\()-] H 1261-73 {1995), ••W. J. Clegg, K. Kendall, N. M, Alford. D. BirL-hall. and T. W. Button, "A Simple Way to Make Tough Ceramics.'" Mwi/rc tLondon). Ml 110] 445-47 (1990), ^L, Zhang and V. D, Krstit. "High Toughness Silicon Carhidc/Graphite Laminar Composite by Slip Casting." Theor. Appl. Fract. Mech.. 24. 13-19 (1995), "A. J, Phiilipps. W. J. Clegg, and T. W. Clyne. "Fracture Behavior of Ceramic Laminates in Bending: 11. Comparison of Mi>deling with Expenniental Data." Acia Meiall- Muier. 41 13| 819-27 |1993). ''T. Ohji. Y. Shigegaki. T. Miyajima, and S, Kan^aki, "Frauture Resistance Behavior of Multilayercd Silicon Nitride," / Am. Cenim. .Soc.. 80 [4] 991-94 (1997), '"T, Chanier. D. Merle, and J. L. Besson. "Laminar Ceramic Composites." J. Eur. Ceram. Soc. IS. 101-107(1995). "Y. Huang. H. N. Hao, Y. L. Cben. and B. N, Zbou. "Design and Preparation of Silicon Nitride Composite with High Fracture Toughness and Nacn." Structuri.'." Acta Meiall. Sin.. 9 [6] 479-84 (19%)- ''H . Liu and S. M. Hsu, "Fraciure Behavior of Multilayer Silicon Nilridc/Boron Nitride Ceramics." / Am. Ceram. Soc, 79 [9| 2452-57 (1996). ' "E. H. Lutz and M. V. Swain, "Fracture Tougbness and Thermal Sbock Behavior of Silicon Niiride-Boron Nitride Ceramics." J. Am. Ceram. Soc, 75 [ I) 67-70 {1992). '"D. Kovar. M. D. Thouless. and J. W, Halkiran. "Crack Deflection and Propagation in Layered Silicon Nitride/Boron Nilride Ceramics." J. Am. Ceram. Soc. 81 14| 1004-12 (1998). '^Z. Lences. K. Hirao. M. E. Brito. M. Toriyama, and S, Kan/aki. "Layered Silicon Niiride-Based Composites witb Discontinuous Boron Nitride Inierlayers." / ,4m. Ceram. Soc. 83 |IO| 2503-508 (1998), " T . Obji. Y, Shigegaki, N. Kondo, and Y, Suzuki, "Fracture Toughness of Multilayer Silicon Nitride with Crack Dellection." Miller. Uii.. 40, 2S0-84 11999). "C. A.Wang, Y. Huang, and Z. P. Xie, "The Effect of Whisker Orientation in SiC Whisker-Reinforced Si,N4 Ceramic Matrix Composites."/ Eur. Ceram. Soc. 19 [10] 190.3-909 (1999). '"C. A, Wang. Y. Huang, and Z. P, Xie. "Improved Resistance to Damage of Silicon Carbide-Wbisker-Reinforced Silicon Nitride-Malrix Composites by WhiskerOriented Alignmeni," i. Am. Ceram. Soc. 84 |1] !fil-64 (2001). '''Y, Huang, H. Guo. and Z. P. Xie, "The Fine Micrtv-Struclure of Interface Layer for Laminated SiiNj Ceramics," J. Maler. Sci. U'li.. 17. 569-7! (1998). ""S. V. Cai, "Study on Biomimetic Design. Preparation and Mechanical Properties of Laminated Si^Nj Matrix Ceramics"; HUJ. Dissertatiop, Tsinghua University. Beijing, People's Republic of China, 1998. D
An Ceram soc851012462-68(2002 Effect of ph of Medium on Hydrothermal Synthesis of Nanocrystalline Cerium(IV)Oxide Powders Nan-Chun Wu, Er-Wei Shi, Yan-Qing Zheng, and Wen-Jun Li Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 201800, People's Republic of China Well-crystallized cerium(IV) oxide( CeO,) powders with nano- Cerium hydroxide formed from Ce(Iv) or Ce(lll) salt solutions sizes without agglomeration have been synthesized by a hydro- n be used as a precursor for hydrothermal synthesis of CeO, hermal method in an acidie medium by using cerium hydrox- powders. A series of precursor reactions such as polymerizing ide gel as a precursor. The relationship between the grain size, dehydrating, and dehydrogenating will occur under hydrothermal the morphology of the Ceo, crystallites, and the reaction conditions and play a key role in the formation of CeO, crystal conditions such as temperature, time, and acidity of the lites. Otherwise, acidity of the hydrothermal medium has direct medium was studied. The experiments showed that with influence on the structure and reaction of cerium hydroxide increasing reaction temperature and time, the CeO, crystal. Therefore, it is very tant to understand the formatio lites grew larger. The crystallites synthesized in an acidic mechanism of CeO, crystallites in the hydrothermal reaction hydrothermal medium were larger and had a more regular mediums with different acidities. In this article, the effect of the morphology than the ones synthesized in a neutral or alkaline pH of the reaction medium on the crystallization of CeO, grains medium when the reaction temperature and time were fixed. under hydrothermal conditions when cerium hydroxide is used as The CeOz crystallites synthesized in an acidic medium were a precursor is reported, as well as the synthesis of nanocrystalline monodispersed; however, there was vigorous agglomeration CeO, powders without agglomeration. among the grains synthesized in a neutral or alkaline medium. It was demonstrated that the hydrothermal treatment was an Ostwald ripening process and the acidity (pH) of the used IL. Experimental Procedure hydrothermal medium played a key role in the dissolution of (1) Hydrothermal Synthesis of Ceo, Powders smaller grains. It is proposed that the dissolution process can The starting materials were cerium(IV) sulfate tetrahydrate control the kinetics of the growth of larger grains. Ce(SO)4H,O)and sodium hydroxide(NaoH)(AR grade purity) The appropriate amounts of Ce(SO4)*4H,O and NaoH were rest Introduction tively dissolved in distilled water to form 0. IM Ce(SO,), and O4M NaOH solutions. When the Ce(SO,), solution was added to the C ERIA(CeO,) is currently being used not only as an oxygen ion NaOH solution, precipitates(gel) were formed immediately. After conductor in solid oxide fuel cells(SOFCs)and oxygen being filtered and washed by distilled water, the precipitates (gel) as monitors'but also as catalytic supports+- of automotive exhaust a precursor were put into an autoclave with a reaction chamber of 40 systems because of its high oxygen ion conductivity. It has been cm. Three quarters of the volume of the chamber was filled with considered that the properties of the materials may be greatly solutions. The acidity of the reaction medium was adjusted by adding mproved if ultrafine powders are used as a raw material. There HCI or NaOH solutions. The hydrothermal experiments were per- fore, preparation of ultrafine CeO, powders without agglomeration formed in the media with pH values of 2. 7 and 12. The hydrothermal has been intensively investigated theses were carried out in a temperature range from room Several techniques that include hydrothermal synthesis, -l urea- temperature to 245C and the time ranged from 18 to 120 h based homogeneous precipitation, hexamethylenetetramine corresponding to a pressure range from I to 60 bar. After the based homogeneous precipitation, coprecipitation, decomposition hydrothermal reaction, the products(powders) were filtered, washed of oxalate precursors, and mechanical mixing have been devel with distilled water several times, and air-dried at l10 C for 7 h ped for the production of ceria or cation-doped ceria particles. As a low-temperature and wet-chemical technique, hydrothermal methods (2) Characterization of the Powders offer an exciting possibility for the synthesis of high-purity, homoge The phase identification of the products was carried out by X-ray neous, and ultrafine powders. Zhou and Rahaman" reported the powder diffraction(XRD; Rigaku RAX-10) Grain sizes were esti- sintering behaviors of CeO, and Y, O,doped CeO, powders synthe- mated by line-broadening analysis according to the formula D sized by a hydrothermal method from Ce(NO, ), solution and ammo- 0.9M(B cos e), where D is the grain size, A is the wavelength of nium hydroxide solution. Tani and co-workers"studied the effect of - rays, e is the diffraction angle, and B is the half-width of the sed mineralizers on hydrothermal synthesis of CeO, powders diffraction peaks. The morphology of the crystallites was observed by Hirano and his co-workers showed controlled crystallization of transmission electron microscopy (TEM, JEOL JEM-2010).The CeO, particles during hydrothermal synthesis from Ce(IV) salts. It chemical composition of the products was determined by X-ray has been found that the agglomeration of hydrothermal CeO, powders fluorescence spectroscopy (XRFS, Philips TW-2404)and inductively is vigorous though the particle size is very fine. coupled plasma spectroscopy (ICP, Perkin-Elmer PLASMA-2000) The products were also characterized by infrared spectroscopy (IR, Perkin-Elmer PE-1600), thermogravimetry, and differential thermal analysis (TG-DTA, Netzsch STA-429) 1-w. Chen-contributing editor Il. Results () Phase and Grain Size mpt No. 188822 Received January 10 The analysis results for the chemical composition of the used upported by the National Sciences Foundation of China precursor and the hydrothermal powders synthesized in the reac. tion medium with different pH values at 200.C for 18 h are shown 2462
journal (2(K)2| Effect of pH of Medium on Hydrothermal Synthesis of Nanocrystalline Cerium(IV) Oxide Powders Nan-Chun Wu, Er-Wei Shi, Yan-Qing Zheng, and Wen-Jun Li Shanghai Institute of Ceramics. Chinese Academy of Sciences, Shanghai 201800, People's Republic of China Well-crystallized cerhim(IV) oxide (CeOj) powders with nanosi7.es without ag^lomenUion have been synthesized by a hydroIhernial method in an acidic medium by using cfriuni hydroxide ^el as a precursor. The relationship between the yrain size, the morphology of the CeO, crystallites, and the reaction conditions such as temperature, time, and acidity nf the medium wa.s .studied. The experiments showed that with increasing reaction temperature and time, the CeO^ crystallites grew larger. The crystallites synthesized in an acidic hydrothermal medium were larger and had a more regular morphology than the ones ,synthesized in a neutral or alkaline medium when the reaction temperature and time were fixed. The CeO, crystallites synthesized in an acidic medium were monodispersed; however, there was vigorous agglomeration among the grains ,synthesized in a neutral or alkaline medium. It was demonstrated that the hydrothermal treatment was an Ostwald ripening process and the aeidity (pH) of the used hydrothermal medium played a key role in the dissolution of smaller grains. It is propo.sed that the dissolution process can control the kinetics of the growth of larger grains. I. Introduction C ERIA (CeOj) is currently being used not only as an oxygen ion conductor in solid oxide fuel cells (SOFCs)''^ and oxygen monitors"' but also as catalytic supports'*"'' of automotive exhaust sysiem.s because of its high oxygen ion condLictivity. It has been considered ihiit the properties of the materials may be greatly improved if iillrafine powders arc used as a raw material. Theretore, preparation of ullral'ine CeO, powders without agglomeration has beeti intensively investigaled. Several techniques that include hydrothermal synthesis/"" ureabased homogeneous precipitation, hexamethy lenetetrami nebased homogeneous precipitation, coprecipitation. decomposition of oxalate precursors."' and mechanical mixing'^ have been developed for the pmduction o\' ceria or cation-doped ceria particles. As a low-temperature and wct-chemical technique, hydrothermal methods olTer an exciting possibility tor the synthesis of high-puHiy. homogeneous, and ultrafine powders. Zhou ;ind Riihaman'' reported the sintering behaviors of CeO-, and Y,O,-doped CcOi powders synthesized by a hydrothermal method Irom Ce(NOi), solution and ammonium hydroxide solution. Tani and ctvworkers'" studied the effect of used tiiincralizers on hydrothermal synthesis of CeO, powders. Hiram) and his co-workers" showed controlled crystallization of CeO-. particles during hydrothermal synthesis from Ce(lV) salts. It has been luund that ihe agglomeration of hydrothennal CeO, powders is vigorous though the particle size is veiy fine. l-W. Chen—contributing editor Manuscript No. 188822. Received January 10, 2O(K1; approved March ]X 2(HI2, .Supported by the National Sciences Foundation of China (NSFC) under Grant No. Cerium hydroxide formed from Ce(IV) or Ce(IIl} sail solutions can be used as a precursor for hydrothermal synthesis of CeO-, powders. A series of precursor reactions such as polymerizing, dehydrating, and dehydrogenating will occur under hydrothermal conditions and play a key role in ihe formation of CeO; crystallites. Otherwise, acidity of the hydrolhermal medium has direct influence on the structure and reaction of cerium hydroxide. Therefore, it is very important lo itnderstand the foniiation mechanism of CeO, crystallites in the hydrolhermal reaction mediums with different acidities. In this article, the effect of the pH of the reaction medium on the crystallization of CeO, grains under hydrothermal conditions when cerium hydroxide is used as a precursor is reported, as well a.s the synthesis of nanocrystalline CeO, powders without agglomeration. II. Experimental Procedure (1) Hydrothermal Synthesis of CeO, Powders The starting materials were cerium(IV) sulfate tetrahydrate (Ce(S04),-4H,0) and sodium hydroxide (NaOH) (AR grade purity). The appropriate amounts of Ce(SO4),-4H2O and NaOH were respectively dissolved in distilled water loIbnn O.\M CeiSO^), and 0.4M NaOH solutions. When the Ce(S04), solution was added to the NaOH solution, precipitates (gel) were fomied immediately. After being filtered and washed by distilled water, the precipitates (gel) as a precursor were put into an aul(x:lave with a reaction chamber of 40 cm"*. Three quartei"s of the volume of ihe chamber was filled wilh solutions. The acidity of the reaction medium was adjusted by adding HCl or NaOH solutions. The hydrothermal experiments were performed in the media with pH values of 2. 7, and 12, The hydrolhermal syntheses were carried out in a temperature range tmm RK>m temperature to 245X and the time ranged from 18 to 120 h, corresp(>nding tt) a pressure range frotii I to 60 bar. After the hydrothermal reaction, the products (ptiwders) were filtered, washed with distilled water several times, and air-dried at 110°C for 7 h. (2) Characterization of the Powders The phase identification of the pr(Hiucts was carried out by X-ray powder diffraction (XRD: Rigaku RAX-10), Grain sizes were estimated by line-broadening analysis according to the fomiula O = 0.9A/(P cos 0), where D is the grain size, X is ihe wavelcnglh of X-rays. 6 is the diffraction angle, and p is ihe half-width of the diffraction peaks. The morphology of the crystallites was obser\'ed by transmission electron microscopy (TEM. JEOL JEM-2010). The chemical composition of the products was determined by X-ray lluoresccnce spcctroscopy (XRFS. Philips TW-2404) and inductively coupled plasma spectroscopy (ICP, Pcrkin Eltiier PLASMA-2()(X)). The products were also ch;iraclerized by infrared speclroscopy {IR, Perkin-Eimer PE-16(X)). themiogravimctry, and differential thermal analysis (TG-DTA, Netzsch STA-42y). HI. Results (1) Phase and Grain Size The analysis results for the chemical composition of Ihe used precursor and the hydrothermal powders synthesized in the reaction medium with different pH values at 200 ^ for 18 h are shown 2462
October 2002 Hydrothermal Synthesis of Nanocrystalline Cerium(/V)Oxide Powders 2463 Table L. Chemical Compositions of the Used Precursor and the Hydrothermal Powder eaction temp pH value of Component content (wti and time used medium CeO, H O CO, S 40 Precursor 8312.71.53 N 1200°C,8h 12 88.66.81.42 95.34.40.20.4 in Table I It can be found that the lower the ph of the reaction medium, the lower the moisture content of the sample, which means that the dewatering of the precursor in the acid medium we more complete than that in the basic medium. The XRD patterns of the samples are shown in Fig. I. The used precursor and the 10 hydrothermal powders displayed all of the major diffraction peaks of CeO, with the fluorite structu The relationship between reaction temperature and the grain 0 size of the powders synthesized in the hydrothermal medium with different acidities is shown in Fig. 2. Figure 3 shows the relation- 100 200 hip between reaction time and the grain size of the powders synthesized at200°C TEMPERATURE (C) (2) Morphology Fig. 2. Relationship between the particle size of the powders and the reaction temperature. The reaction time was fixed at 18 h.(.) Acidic TEM photographs of the used precursor and the powders medium neutral medium:(A) basic medium. synthesized at different hydrothermal reaction temperature for 18 h are shown in Fig. 4. Some TEM photographs of the powders synthesized at 200C for different reaction times are acidities are displayed in Fig. 6. It is well known that the broad absorption band located in the area from 3200 to 3600 cm - In this paper, it is obvious that hydrothermal treatment of approximately corresponds to the O-H stretching vibration,and Ce(OH), is an Ostwald ripening process and the acidity has he one located in the area from 400 to 750 cm to the Ceo considerably influenced the hydrothermal treatment process. stretching vibration. The absorption peaks at 1600. 1500, or 1340 cm, and 1060 cm correspond to the H,O bending vibration. (3) IR Spectrograms the CO; stretching vibration, and the Ce-OH or SO4 stretching c IR spectrograms of the used precursor and hydrothermal pow vibration, respectively. From Fig. 6. it can be concluded that the synthesized at 200 C for 18 h in the medium with various crystallization of the powders synthesized in an acidic hydrother- mal medium was more complete than others, which agrees with the (111) processes are associated with SO? and Co? othermal t given in analysis for the chemical compositions of the samples Table L. Figure 6 and Table I show that the hydr treatmet (D) 220) (311) (4) TG-DTA Spectrograms The TG-DTA spectrograms of the used precursor and the hydrothermal powders are shown in Fig. 7. The ascent of the TG curve denotes an increase of sample weight, and it is expressed by positive Am: and the descent of the TG curve denotes a decrease (C) 、∧八 且 0 100 20(deg TIME (h) Fig. 1. X-ray diffraction patterns of the used precursor(A)and the Fig. 3. Relationship between the particle size of the powders and the hydrothermal powders synthesized at 200 C for 18 h in the reaction reaction time. The reaction temperature was fixed at 200C(+)Acidic medium with a pH of 12(B), 7(C), and 2(D) medium;() neutral medium: (X) basic medium
October 2{)O2 Hvdrothermal Synthesis of Nanocry.<italline Cerium(lV) O.xide Powders 2463 Tiihtc I. Chemical Compositions of the I sed Precursor and the Hydrothermal Powders u....,.,i .,„ „» f .!,» Componcni L-onreni |wi*} Precursor Powder 1 Powder 2 Powder 3 unJ liiTiL' 200"C. 18 h usL'd mciiiuin 12 7 2 CeO- 83 88.6 90.1 95.3 H.O 12.7 6.8 5.8 4.4 CO, 1.5 1.4 1.1 0.2 so.. 3.2 2.9 2.7 0.4 in Table I. It can be found that the lower the pH of the reaction niedium. the lower the moisture content of the sample, which means ihiit the dewalering of ihe precursor in the acid medium was more complete than thai in the basic medium. The XRD paitems of the samples are shown in Fig. I, The used precursor and the hydrothermal powders displayed all of the major diffraction peaks of CeOi with the fluorite structure. The relationship between reaction temperature and the grain size of Ihe powders synthesized in the hydrothermal medium with different acidities is shown in Fig. 2. Figure 3 shows the relationship between reaction time and Ihe grain size of the powders synthesized at 200"C. (2) Morphology TEM photographs of the used precursor and the powders synthesized at different hydrothermal reaction temperature for 18 h are shown in Fig. 4. Some TEM photographs of the powders synthesized at 200°C for different reaction times are shown in Fig. 5. In this paper, it is obvious that hydrothermal treatment of Ce{OH)4 is an Ostwald ripening process and the acidity has considerably influenced the bydrothermal treatment process. ii) IR Spectrograms IR spectrograms of the used precursor and hydrothermal powders synthesized at 2OO''C for 18 h in the medium with various (111) 5 N So 50 40 30 20 10 0 100 200 300 TEMPERATURE (°C) Fig. 2. Relationship between the particle size of the powders and llic reaction lemperature. The reaction lime was fixed al 18 h. (•) Acidic medium: (•) neulral nieiiium; (A) basic niedium. acidities are displayed in Fig. 6. It is well known ihat ilie broad absorption hand located in the area from 3200 to 3f)(H) cm ' approximately corresponds to the 0-H stretching vibration, and the one located in the area from 400 to 750 cm ' to the CeO, stretching vibrution. The absorption peaks at 1600. 1500. or 1340 cm the vibration, respectively. From Fig. 6. it can be concluded tbat the crystallization of the powders synthesized in an acidic hydroihermal medium was more complete than others, which agrees wilh the analysis for the chemical compositions of the satnples given in Table I. Figure 6 and Table I show that tbe hydrothermat treatment processes are associated witb SOj " and CO^ . ~', and 1060 cm ' correspond to the H,O bending vihralion, COt" stretching vibration, and ihe Ce-OH or SOj stretching (4) TG-DTA Spectrograms The TG-DTA spectrograms of the used precursor and the bydiotbermal powders are shown in Fig. 7. The ascent of the TG curve denotes an increase of sample weight, and it is expressed by positive Am: and the descent of the TG curve denotes a decrease 20 50 100 TIME (h) 150 Fig, 1. X-ray diffraction patterns of the used precursor (A) and the hydrothermal powders synthesized at 200°C for 18 h in the reaction medium wUh a pH of 12 (B). 7 (C), and 2 (D). Fig. 3. Relationship between the particle si7e oi ihe powders and the reaction time. The reaction temperature was fixed al 20O''C. (•) Acidic medium; (•) neuiral medium: (X) basic medium
Journal of the American Ceramic Socieny-Wu er al. Vol. 85. No. 10 (A) (B) e505q己 40 nm 40 nm 505H4q 3b052b8b 40 nm 28 nm (E) 40 nm Fig. 4. TEM photographs of the used precursor (A)and the hydrothermal powders synthesized at 245 C in the neutral medium(B) and the basic medium (C), and at 200(D) and 245 C(E) in the acid medium. The reaction time was fixed at I8 h of sample weight, and it is expressed by negative Am. When the IV. Discussion two parts, the total Am is equal to the algebraic sum of the two Formation o Ceo crystallites Synthesized under arts Am. Similarly, when the DTA curves are out of range and are divided into two parts, the whole curve may be obtained by Because the solubility products of some precipitated sub- joining the two parts after shifting vertically the divided part. stances are bigger (for example, K p(AgCI) The dotted curves indicate a temperature change from room K p(Baso4)=I I X 10,etc). when Ostwald ripening occurs temperature to 1000 C, an ascent of the curve denotes a rise in the smaller grains dissolve more quickly, the solute diffuses more temperature, and a descent of the curve denotes a fall in slowly, and the larger grains grow more slowly: therefore, it is temperature. The whole curve may be obtained by joining the assumed that the growth kinetics are controlled by diffusion or divided parts after shifting vertically the divided parts. The surface reaction. However, the solubility product of Ce(OH)a endothermic peaks with the temperature below 150C are (i.e. CeO, 2H- O)is 2 x 10 and is far smaller than for one of elated to dehydrating of the examined samples those precipitated substances. When Ostwald ripening occurs, the
2464 Jtnirnal of the American Ceramic Societ\—Wii et al. Vol. 85. No. 10 Fig. 4. TEM phoiographs of lhe used precursor (A) and the hydrolhermal powders synthesized al 245''C in the neutral medium (B) and the basic medium (C). and ai 200'' (D) and 245X (E) in the acid medium. The reaction time was fixed at 18 h. of satiiple weight, and it i.s expressed by negative Am. When the TG curves are out of range of the plotter and are divided into two parts, the total A;» is equal to the algebraic sum of the two parts Am. Similarly, when the DTA curves are out of range and arc divided into two parts, the whole ctirve may be obtained by joininj! the two parts after shifting vertically the divided part. The dotted curves indicate a temperature change from room tetiiperauirc to 1000°C. an aseent of the curve denotes a rise in temperature, and a descent of the curve denotes a fall in temperature. The whole curve may be obtained by joining the di\ided parts after shifting vertically the divided parts. The endotherinic peaks with the temperature below ISO^C are related to dehydrating of the examined samples. IV. Discussion (1) Formation of CeOi Crystallites Synthesized under Hydrothermal Conditions Because the solubility products"* of some precipitated substances are bigger (for example. A',,,(AgCI) ^ I.S X 10 '" and A',p(BaSO4) - 1.1 X 10 '". etc), when Ostwald ripening occurs the smaller grains dissolve more quickly, the solute diffuses move slowly, and lhe larger grains grow more slowly; therefore, it is assumed that the growth kinetics are controlled by diffusion or surface reaction. However, the solubility product"* of Ce(0H)4 (i.e., CeO;-2H,O) is 2 X 10 •"* and is far smaller than for one of those precipitated substances. When Ostwald ripening occurs, the
October 2002 Hydrothermal Synthesis of Nanocrystalline Cerium( /V) Oxide Powders 250524q(B) S05449b 40 nm 505242 S054244 (C) (D 40 nm 40 nm Fig. 5. TEM photographs of the hydrothermal powders synthesized in the neutral medium for 72(A)and 120 h(B), and in the acidic medium for 72(C) and 120 h(D). The reaction temperature was fixed at 200'C smaller grains dissolve far more slowly, and in addition, grains of where Au is the driving force of the crystallization, k is the precipitated Ce(OH)a cluster together to form agglomerates, so the Boltzmann constant, and T is the absolute temperature. In the solute diffuses more quickly and the larger grains grow more solution systems Au can be expressed as quickly. Therefore, it is assumed that the growth kinetics are controlled by dissolution △μ= kT In C/c Ce(OH)a is a basic pr [OH I will lead to n apparent decrease of solubility of Ce(OH)a, and increasing where C is the concentration of the supersaturated solution, and Co [HI will lead to a sizable increase of solubility of Ce(OH)a is the concentration of the saturated solution From Eqs. (1)to (3). For Ostwald ripening, there is an equilibrium value for grain the following formula can be obtained size r. When r is smaller than r, smaller grains will dissolve and eventually disappear. When r is larger than r", larger grains will R= B(C-Co)/Co grow. The relationship between grain size and solubility may be expressed by the Gibbs-Thomson equation: formula shows that, with increasing C, R will increase ally; and with decreasing C, R will decrease gradually unti C(r)=C, exp(2yV/vRTr equal to Co and a new dissolution-recrystallization equilib- rium is established. This can be represented as where C(r) is the solubility of a grain with radius r, Ca is the normal equilibrium solubility of the substance (r- oo), y is the 2CeOH=[CeOCe]+H,O (Ks= 16.5) (5) interfacial tension, V is the molar volume of the solute, v is the number of ions in the formula unit, r is the grain radius. R is the In this paper, the concentration of the species CeOH is taken for gas constant, and T is the absolute temperature. the concentration of the solute and is equal to C. In Eq.(4), Cis According to the maximal growth law of crystal growth. the qual to C(r) and Co is equal to C(the solubility while r is equal growth rate of the crystal can be written as to r). So Eq (4)can be represented as R= B[C, exp(2yV/vRTr)/C-11 where R is the growth rate of the crystal. B is a kinetic coefficient. This formula indicates that R is associated with C and o is the relative supersaturation. It is assumed that the maximal Strictly speaking, the formula of the reactant or product in growth law holds true in this paper. o is defined as he hydrothermal medium should be written as Ce, O,(OH)(H,O), (SO4),(CO). However, in this paper, the o=exp(△p/k7 2) reactants or products are respectively represented as Ce
October 2002 Hydrothermal Synthesis of Nanocrystalline Ccrium(lV) Oxide Ponders 2465 Fig. 5. TEM photographs of the hydrothermal powders synthesized in the neutral medium lor 72 (A) and 120 h (B), and in t!ie acidic medium for 72 (C) and 120 h |D). The reaction lemperature was fixed al 200°C. smaller grains dissolve far more slowly, and in addition, grains of precipitated Ce(0H)4 cluster together to form agglomerates, so the solute diffuses more quickly and the larger grains grow more quickly. Therefore, it is assumed that the growth kinetics are controlled by dissolution. CefOHjj is a basic precipitate, so increasing |0H~] will lead to an apparent decrease of solubility of CXOHjj. and increasing [H"*"] will lead to a sizable increase of solubility of CXOHjj. For Ostwald ripening, there is an equilibrium value for grain size r*. When r is smaller than r*. smaller grains will dissolve and eventually disappear. When ;• is larger than /•*. larger grains will grow. The relationship between grain size and solubility may be expressed by the Gibbs-Thomson equation:'*^ C{r) = C. e\p{2yVJvRTr) where C(r) is the solubility of a grain with radius r. C^. is the normal equilibrium solubility of the substance (r —> oc), -y is the interfacial tension. U,,, is the molar volume of tbe solute. \' is the number of ions in the formula unit, r is the grain radius. R Is the gas constant, and T is the absolute temperature. According to the maximal growth law of crystal growth,^" the growth rate of the crystal can be written as R-^u , (1) where R is the growth rate of the crystal. ^ is a kinetic coefficient, and (T is the relative supersaturation. It is assumed that the maximal growth law holds true in this paper, a is defined as where Afi is the driving force of the crystallization, k is the Boltzniann constant, and T is the absolute temperature. In the solution systems Ap. can be expressed as = kT In C/Q, (3) where C is the concentration of the supersaturated solution, and C(, is the concentration of the saturated solution. Prom Eqs. (I) to (3). the following formula can be obtained: (4) This formula shows that, with increasing C, R will increase gradually; and with decreasing C R will decrease gradually until C is equal to C,, and a new dissolution-recrystallization equilibrium is established. This can be represented as"' = [CeOCe]"' + 16.5) (5) In this paper, the concentration of the species CeOH"''^ is taken for the concentration of the solute and is equal to C,. In Eq. (4). C is equal to C{r] and C^, is equal to C* (the solubility while r is equal to /•*). So Eq. (4) can be represented as R = exp{2yVJvRTr)/C* - I] - 1 (2) This formula indicates that R is associated with C^, Strictly speaking, the formula of the reactant or product in the hydrothermal medium should be written as Ce,,O,(0H),,(H.0),(SO4),(C0,)- However, in this paper, the reactants or products are respectively represented as Ce''
Joumal of the American Ceramic Sociery-Wu et al. Vo.85.No.10 Equations(10).(11), and(12) can be obtained respectively from Eqs. (6),(7), and (8)combined with Eq (9). When the concentra tion of CeOH increases, other reactions take place 50 CeOh'+HCo:+OH- H,O CeOHCO,+ H,O (Kn=9.02X 10) 8az (C) CeCo,+ H (13 Ce-OH or SO.2 2CeOH-[CeOCe]*+H,O (Ks= 16.5) HO CO2- Under certain conditions, the following reactions may take place: 8 OH CO, HCO5=H+CO3(K4=4.7×10-1) (14) 0 (B) Ce-OH CeOH++H= Ce"+H,O (KIs=0. 192) or SOa Reaction (15) is the backward reaction of reaction(8). In Ref. 23 50 H,O CO3 the following reaction is proposed CO 2 Ce++NO3+HO=Ce(OH)(NO3)2++H+(K16=102) In Refs. 24 and 25. some evidence otherwise shows that Ce(O Ce-OH H)(CO,)is stable, so it is suggested that the following reaction can H,ODor SO2 occur OH CO, 2-CO.2- Ce"+CO:+H,0-CeOHCO, +H(Kn,= 10) 40003000 2000 CoCO,+ H In fact, because CO, has two negative charges and will produce WAVENUMBER (cm") greater crystal lattice energy than NO,, we think that a Ce complex with CO is more stable than a Ce complex with NO3. Moreover, it appears that CeOCO, can be stable as Fig. 6. IR spectra of the used precursor (curve A)and the hydrothermal Ce(OH)(CO, ). By combining Eqs. (9).(14).(15), and(17),E powders synthesized in the basic medium (curve B), the neutral medium(13) can be obtained, as shown above (curve C), and the acidic medium(curve D). The reaction temperature and the time were fixed at 200C and 18 h Hydrolysis and dimerization reactions took place successively (Eqs.(8)and (5)), leading to the formation of CeO, grai Because the concentration of CeOH increased rapidly, a great number of CeO, nuclei formed in a very short time, equations CeOH+, CeSO2*, CeOHCO', etc, in each chemical equation (10). (ID), and(3)show that in basic solution CeSO4, Ce(Soa)e, for simplicity. In this paper, the equilibrium constants at 25C re applied. The effect of the temperature on the equilibrium complex with SO: or CO to be resolved, so that it is difficult constant of the reaction limited only to Eqs. (5)and(8)are for the nuclei which combine with SO or CO on the surface discussed: the effect of the temperature on other reactions is not to grow further. discussed for lack of thermodynamic data. Under certain Under certain conditions. the following equilibria can be estab conditions, all reactions discussed in this paper are reversible lished: In Ce(SO1), solution some chemical equilibria are estab- Ce(SO)+H'= CeSO"+HSO, (KIs=5X 10) lished:22-21 (18) Ceso:+ HSO:=Ce(SO)+H(K-200 CeSO+H'=Ce4+HSO(K19=286×10)(19) Ce"+ HSO:= CeSO4+H(K,=3500 CeOHCO,+H= Ce+CO:+H,O(K0=10) Ce+H O= CeOH'+H(K=5.2) 8) Reactions(18),(19), and (20)are the backward reactions of egs (6),(7), and (17), respectively. With the sum of reactions(18) H+OH=H,O(K,=10 4) (19) and(8), an overall reaction can be obtained Ce(SO4)+H+HO- CeOH++2HSON When the Ce(SO4), solution was added to the NaOH solution to (K21=7.42×10) (21) prepare precursor, the following reactions took place CeSO:+OH"+ HSO: =Ce(SO)2+H,O With the sum of reactions(8) and (20) and the backward reaction of Eq.(14), an overall reaction can be obtained (K10=200×10) CeOHCO3+H=CeOH+HCO3(K2=1.1×10° Ce++OH"+ HSO.=CeSOf'+HO (K1=3500×10 (I) Under hydrothermal conditions, with increasing temperature the quilibrium of reaction (8)shifts toward the right, so that the concentration of the hydrogen ion is increased. This leads to a Ce+OH- CeOH (K=5.2 X 10) (12) transformation such that the equilibria of reactions(21)and(22)
2466 Journal of the American Ceramic Society—Wu et al. Vol. 85. No. 10 i 0 4000 3000 2000 1000 WAVENUMBER (cm Fig. 6. IR spectra of [lie used precursor (curve A) and [he hydrothermal powders syntht'si/.eij in the hasic medium (curve B). ihe neuirai medium (curve C). iinJ ihe acidic medium (curve D). The reaction temperature and the lime were fixed at 200"C and 18 h. CeOH '. CeSO^^, CeOHCOt. etc., in each chemical equation lor simplicity. In this paper, the equilibrium constants at 25''C are applied. The effect of the temperature on the equilibrium constant of the reaction limited only to Eqs. (5) and (8) are discussed; tbe effect of the temperature on other reactions is not discussed for lack of thermodynamic data. Under certain conditions, all reactions discussed in tbis paper are reversible. In CefSOj)-, solution some chemical equilibria are establisbed:--'''" ^ HSO; = J: + H' = 200) + HSO4 = CeSOf + H' [Ky = 3500) Ce^' + H:0 H* + OH" - H.O + H" {K^ = 5.2) - 10'*) (6) (7) (8) (9) When the Ce(S04)2 solution was added to the NaOH solution to prepare precursor, the following reactions took place: CeSOf + OH" + HSO; - CelSOJ; + (/Tu, = 200 X lO'-*) e'^ + OH + HSO4 - H.O 10' OH - CeOH^" (^,2 - 5.2 x (10) (II) (12) Equations (10), (II), and (12) can be obtained respectively from Eqs. (6). (7). and (8) combined witb Eq. (9). When the concentration of CeOH"'"" increases, other reactions take place: CeOH'' +HCO," +0 H = CeOHCo; + H2O (A:,, = 9.02 x 10") U CeOCO., + H' (13) 2CeOH-^ = [CeOCe]*^^ + H2O {K, = 16.5) (5) Under certain conditions, the following reactions may take place:'** HCO3 - H* + CO^" (A:|4 - 4.7 X 10"") (14) CeOH'' + H" =Ce-" + H2O (/f,, = 0.192) (15) Reaction (15) is the backward reaction of reaction (8). In Ref. 23 the following reaction is proposed: Ce-*" + NO3" + H2O = Ce(0H)(N03)^^ + H* (A'.fi = 10') (16) In Refs. 24 and 25. some evidence otherwise shows that Ce(OH)(CO,) is stable, so it is suggested that tbe following reaction can occur: + col- + H2O - CeOHCO," + u CeOCO, + H' (17) In fact, because CO3 has two negative charges and will produce greater crystal lattice energy than NO^". we think that a Ce"^ complex with CO3" is more stable than a Ce"" complex with NO,^. Moreover, it appears that CeOCO, can be stable as Ce(bH)(CO,). By combining Eqs. (9). (14). (15). and (17). Eq. (13) can be obtained, as shown above. Hydrolysis and dimerizatJon reactions took place successively (Eqs. (8) and (5)). leading to the formation of CeO^ grains. Because the concentration of CeOH'^ increased rapidly, a great number of CeO^ nuclei formed in a very short time. Equations (10), (I I). and (1*3) show that in basic solution CeSO'^, Ce(SO4)2. and CeOHCO,' are quite stable, and It is quite difficult for a Ce"*"^ complex with SO]" or CO^ to be resolved, so that it is difficult for the nuclei which combine witb SO^ or CO^ on the surface to grow further. Under certain conditions, tbe following equilibria can be established: 4)2 + H' - CeSOj' + HSO; = 5 X 10') (18) CeS(^^ + H* = Ce""" + HSO4' (^T,, = 2.86 x 10'') (19) CeOHCO; + H^ - Ce"^ + COf + H.O (Z^,,, = 10"') (20) Reactions (18), (19). and (20) are the backward reactions of Eqs. (6), (7), and (17), respectively. With the sum of reactions (18), (19), and (8), an overall reaction can be obtained: 4)2 + H^ + H.O = CeOH'' + 2HS0; (/r,, = 7.42 X 10-") (21) With the sum of reactions (8) and (20) and the backward reaction of Eq. (14), an overall reaction can be obtained: CeOHCOl + H" - CeOH'^ + HCO," (K22= 10") (22) Under hydrothermal conditions, with increasing temperature the equilibrium of reaction (8) sbifts toward the rights' so that the concentration of the hydrogen ion is increased. This leads to a transformation such that the equilibria of reactit)ns (21) and (22)
