《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_Effect of Cubic Phase on the Kinetics of

Printed in Grcat Britain. All righ PII:S02728842(96)00074-0 Effect of cubic Phase on the Kinetics of the Isothermal Tetragonal to Monoclinic Transformation in ZrO2 3 mol%Y203 Ceramics W.Z. Zhu Institute for Nonmetallic Materials, Swiss Federal Institute of Technology(ETH), CH-8092, Zurich, Switzerland (Received 12 December 1995; accepted 23 July 1996) Abstract: In this paper, the effect of cubic phase on the kinetics of the isothermal gonal (t) to monoclinic(m) transformation in ZrO2(3 mol%Y2O3)ceramics was studied by means of thermal expansion analysis (TEA), scanning electron microscopy(SEM), transmission electron microscopy (TEM)and X-ray diffrac- (XRD). Experiment t-m transformation could be expressed using the Johson-Mehh-Avrami equa tion. 1.e where f refers to the volume fraction of transformed m-phase, k is a variable associated with the energy barrier for critical nucleation and growth, and n is a constant depending on the nucleation sites. It was found that both the"nose empcrature and the incubation periods of the tTt curve of the Zro2( mo 1%Y2O3)ceramics were decreased in comparison with those of the TTt curve of the ZrO2(2 mol%Y2O3) ceramics. The time-temperature-transformation (Ttt) curve is C-shaped, with the"nose"temperature determined to be 300 C. Activa tion energy for transformation was regressed to be kJ/mol, which didn't cubic phase grains refined t-phase grains during hot-pressing. It is proposed that the nucleation and longitudinal growth of m-phase plates in Zro2(Y2O3)cera mics are displacive, while the sidewise growth thereafter is controlled by short range diffusion of oxygen ions and, in this sense, the t-m transformation in ZrO2(Y2O3) ceramics possesses both displacive and non- displacive features. 997 Elsevier Science Limited and Techna S.r.I 1 INTRODUCTION of stabilizer, grain size, shape and site of the t phase grains, as well as constraints imposed by the Transformation toughening through the tetragonal neighbouring matrix, -3many problems still (t)to monoclinic(m) transformation in Zro2 alloys remain unsolved. The stress-triggered t-m transi- is believed to be one of the most effective ways of tion occurring in Zro2 alloys is usually considered improving the toughness and reliability of struc- to be athermal, but may also show some isothermal tural ceramics. Although our knowledge of this character 4-7 This isothermal component of the aspect has advanced considerably, e.g. it is gener transformation has received relatively little atten ally agreed that control of the metastability of the tion. There are divided views on the nature of the t-phase is the key to obtaining optimum toughen- kinetics of the isothermal t-m transformation in ing effect and is affected by such factors as content ZrO2(,) ceramics. Results by Nakanishi and Shi su have shown that the kinetics of the Present address: Department of Materials Science and Engi- isothermal t-m transformation is controlled by neering, Zhejiang University, Hangzhou 310027, P R China. the diffusion of oxygen ions with the implication of

Ceramics Iniernational24 (1998) 3543 0 1997 Elsevier Science Limited and Techna S.r.1. Printed in Great Britain. All rights reserved PII:SO272-8842(96)00074-O 0272-8842/97 $17.00+.00 Effect of Cubic Phase on the Kinetics of the Isothermal Tetragonal to Monoclinic Transformation in Zr02(3 mol%Y203) Ceramics W. 2. Zhu” Institute for Nonmetallic Materials, Swiss Federal Institute of Technology (ETH), CH-8092, Zurich, Switzerland (Received 12 December 1995; accepted 23 July 1996) Abstract: In this paper, the effect of cubic phase on the kinetics of the isothermal tetragonal (t) to monoclinic (m) transformation in ZrOz(3 mol% YzOs) ceramics was studied by means of thermal expansion analysis (TEA), scanning electron microscopy (SEM), transmission electron microscopy (TEM) and X-ray diffrac￾tometry (XRD). Experimental results showed that the kinetics of the isothermal t--tm transformation could be expressed using the Johson-Mehl-Avrami equa￾tion, i.e. f = 1 - exp(-kt”) where f refers to the volume fraction of transformed m-phase, k is a variable associated with the energy barrier for critical nucleation and growth, and n is a constant depending on the nucleation sites. It was found that both the “nose” temperature and the incubation periods of the TTT curve of the Zr02(3mo- 1% Yz03) ceramics were decreased in comparison with those of the TTT curve of the Zr02(2 mol% YzOs) ceramics. The tim&emperature-transformation (TTT) curve is C-shaped, with the “nose” temperature determined to be 300°C. Activa￾tion energy for transformation was regressed to be 22.74 kJ/mol, which didn’t change with temperature. TEM observation revealed that preferential growth of cubic phase grains refined t-phase grains during hot-pressing. It is proposed that the nucleation and longitudinal growth of m-phase plates in ZrOz(YzOs) cera￾mics are displacive, while the sidewise growth thereafter is controlled by short￾range diffusion of oxygen ions and, in this sense, the t--+m transformation in ZrOz(YzOs) ceramics possesses both displacive and non-displacive features. 0 1997 Elsevier Science Limited and Techna S.r.1. 1 INTRODUCTION Transformation toughening through the tetragonal (t) to monoclinic (m) transformation in Zr02 alloys is believed to be one of the most effective ways of improving the toughness and reliability of struc￾tural ceramics. Although our knowledge of this aspect has advanced considerably, e.g. it is gener￾ally agreed that control of the metastability of the t-phase is the key to obtaining optimum toughen￾ing effect and is affected by such factors as content *Present address: Department of Materials Science and Engi￾neering, Zhejiang University, Hangzhou 310027, P. R. China. of stabilizer, grain size, shape and site of the t￾phase grains, as well as constraints imposed by the neighbouring matrix,1-3 many problems still remain unsolved. The stress-triggered t+m transi￾tion occurring in ZrOz alloys is usually considered to be athermal, but may also show some isothermal character.47 This isothermal component of the transformation has received relatively little atten￾tion. There are divided views on the nature of the kinetics of the isothermal t+m transformation in Zr02(Y20s) ceramics. Results by Nakanishi and Shigematsus have shown that the kinetics of the isothermal t-m transformation is controlled by the diffusion of oxygen ions with the implication of 35

W.Z. Zhu bainite transformation, which is quite different where Im(111)and Im(1l1)refer to the relative peak from results of their previous studies. Heuer et strengths of the (111)and(111)planes of the al. pointed out that the t-m transformation in monoclinic phase, respectively, and Ic(lll)repre Y-TZP ceramics displays both athermal and iso- sents the relative peak strength of the (111) plane thermal kinetics; they are both stress-assisted mar- of the tetragonal or cubic pha ple fo tensitic transformations. They viewed the athermal TEM inspection was prepared by the conventional character usually found in a cooling transforma- procedure of grinding, dimpling and ion-thinning tion as due to rapid isothermal transformation of to perforation, followed by coating with a thin film the most easily nucleated regions at a succession of of amorphous carbon to avoid charging and Temperatures. In other words, a range of activation JEM100 type microscopy was operated with energies exists for such martensitic transforma- 120 v as the accelerating voltage I data has been obtained to confirm this deduction. The intent of the present paper is to interpret the character of the 3 RESULTS AND DISCUSSION kinetics of the t-m transformation in ZrO2 (3 mol%Y2O3) ceramics, which consist of cubic, tet ragonal and monoclinic phases, and to propose an 3. 1 Thermal expansion curves of ZrO2 (2 mo/%y203 appropriate mechanism and ZrO2 3 mo/%Y2O3)specimens Figure I shows the thermal expansion curves of 2 EXPERIMENTAL PROCEDURES ZrO2(2 mol%Y2O3) and ZrO2(3 mol%Y2O3) ceramics obtained at a heating and cooling rate of Powders of zirconia containing 2 mol% and 10 C/min. Specimen size initially increases with 3 mol% yttria in solid solution, respectively, with increase in temperature, then drastically decreases purity higher than 99% and average diameter of as the temperature is increased above the a, point about 0-1 um, were cold-pressed into pellets of (starting temperature of the mt transformation), 24 mmx6mmx6mm in size under a pressure of finally it increases again as the temperature is fur- 300 MPa, followed by pressureless sintering at ther increased above the Af point (ending tem 1 600C with 5 h holding time. a relative density of perature of the m+t transformation). It is more than 98% could be obtained. Thermal apparent that at temperatures above the Ar point expansion analysis experiments were performed on only t-phase exists in the specimen. Similarly, the a Perkin-Elmer 7 Series Thermal Analysis System specimen size initially decreases linearly with with polished specimens, machined to be decrease in temperature until the Ms point(starting 4 mmx4mmx3 mm in size. The point at which a temperature of the t-m transformation), below tangential line deviates from the thermal expansion which it increases significantly due to the dilata curve was defined as the transformation tempera- tional nature of the t-m transformation. When ture. The relative amount of monclinic phase was the temperature is further decreased below the m calculated using the formula proposed by Garvie, point(ending temperature of the t-m transfor according to the X-ray diffraction results mation), thermal expansion curves reverse to the linear shrinkage section until ambient temperature (l11+lm(1l It should be noted that, in the present study, Im(111)+Im(111)+lct(111) ZrO2 (2 mol%Y2O3) ceramics were sintered the single t-phase region according to the phas .Oh 00500600700 00500600700 Fig. 1. Thermal expansion curves of ZrO2(Y203)ceramics obtained at a heating and cooling rate of 10C/min (a)ZrO2(2 mol%Y2O3);(b)Zro2(3 mol% Y 2O3)

36 bainite transformation, which is quite different from results of their previous studies.9 Heuer et aLlo pointed out that the t-m transformation in 3Y-TZP ceramics displays both athermal and iso￾thermal kinetics; they are both stress-assisted mar￾tensitic transformations. They viewed the athermal character usually found in a cooling transforma￾tion as due to rapid isothermal transformation of the most easily nucleated regions at a succession of temperatures. In other words, a range of activation energies exists for such martensitic transforma￾tions. Up to now, no experimental data has been obtained to confirm this deduction. The intent of the present paper is to interpret the character of the kinetics of the t-m transformation in ZrOz(3 mol% Y203) ceramics, which consist of cubic, tet￾ragonal and monoclinic phases, and to propose an appropriate mechanism, 2 EXPERIMENTAL PROCEDURES Powders of zirconia containing 2mol% and 3mol% yttria in solid solution, respectively, with purity higher than 99% and average diameter of about O-1 grn, were cold-pressed into pellets of 24 mmx6mm x6 mm in size under a pressure of 300MPa, followed by pressureless sintering at 1600°C with 5 h holding time. A relative density of more than 98% couid be obtained. Thermal expansion analysis experiments were performed on a Perkin-Elmer 7 Series Thermal Analysis System with polished specimens, machined to be 4 mm x4 mmx 3 mm in size. The point at which a tangential line deviates from the thermal expansion curve was defined as the transformation tempera￾ture. The relative amount of monclinic phase was calculated using the formula proposed by Garvie,’ ’ according to the X-ray diffraction results: M% = I,(lii) +~,(lil) I,(lli) +1nl(lll) +Lt(lll) -0.6 - I I I I IT I 0 100 200 300 400 500 600 700 Temperature (“C) W. Z. Zhu where I,( 11 i) and Im( 111) refer to the relative peak strengths of the (11 i) and (111) planes of the monoclinic phase, respectively, and I& 111) repre￾sents the relative peak strength of the (111) plane of the tetragonal or cubic phase. A sample for TEM inspection was prepared by the conventional procedure of grinding, dimpling and ion-thinning to perforation, followed by coating with a thin film of amorphous carbon to avoid charging and JEMlOO type microscopy was operated with 120 kV as the accelerating voltage. 3 RESULTS AND DISCUSSION 3. I Thermalexpansion curves of ZrQ(2 moi% Y&) and ZrUz(3 mol% Y203) specimens Figure 1 shows the thermal expansion curves of ZrOz(2 mol% Y203) and ZrOz(3 mol% Y203) ceramics obtained at a heating and cooling rate of lO”C/min. Specimen size initially increases with increase in temperature, then drastically decreases as the temperature is increased above the A, point (starting temperature of the m-+t transformation), finally it increases again as the temperature is fur￾ther increased above the Af point (ending tem￾perature of the m-t transformation). It is apparent that at temperatures above the Ar point, only t-phase exists in the specimen. Similarly, the specimen size initially decreases linearly with decrease in temperature until the M, point (starting temperature of the t+m transformation), below which it increases significantly due to the dilata￾tional nature of the t-+m transformation. When the temperature is further decreased below the Mf point (ending temperature of the t-+m transfor￾mation), thermal expansion curves reverse to the linear shrinkage section until ambient temperature. It should be noted that, in the present study, Zr02(2 mol% Y203) ceramics were sintered in the single t-phase region according to the phase g 0.2 g ‘;: 0 3 a -0.2 - 5 6h -0.4- .s z z -0.6 - p! I I MS I I I _ 0 100 200 300 400 500 600 700 Temperature (“Cl Fig. 1. Thermal expansion curves of Zr02(Y203) ceramics obtained at a heating and cooling rate of lO”C/min: (a) Zr02(2mol% Y,O,); (b) Zr02(3 mol% Y20,)

Isothermal tetragonal to monoclinic transformation Table 1. Starting and ending temperatures of the t-+m and m-t transitions for ZrO2(3 mol%Y2O3) and zrO2(2 mol%Y2O3) specimens sintered at 1600c M A ZrOz (3 mol%Y2O3) 387°C 142°C 374°C 542°C ZrO2(2 mol%Y2O3 459C 322°C 567°C 602C 0.8(a) 08(b) 85945四 Fig. 2. Thermal expansion curves obtained at a heating and cooling rate 100 C/min for ZrO2 (3 mol%Y2O3) ceramics held at Table 2. Results of thermal expansion and XRD measurements for ZrO2(3 mol%Y2O3)specimens held at differ ent temperatures for 4 h Temperature(°C) 200 e-m phase 13.4% sothermal m-phase 52.47% Athermal m-phase Overall m-phase Expansion before hol 3532% 0.24335 Expansion during holding n after holding 24 00589%04098% Overall expansion 1.3990% 1.1510% 15120% diagram. 2 The microstructure of the t+m dual is evident that a cooling rate of 100 C/min is not phasc obtaincd at room tempcrature implies that high enough to inhibit the t-m transformation part of the sintered t-phase has transformed to m- before g, and phase during cooling. However, ZrO2 (3 mo- perature, the more pre-Im phase produced prior Lo 1%Y2O3) ceramics were sintered in the c+t dual holding. Therefore, the initial phase constitutent phase region and the microstructure obtained at for the measurement of the isothermal kinetics room temperature is composed of c, t and m-pha ses. Table 1 lists the transformation temperatures of ZrO2(2 mol% Y2O3)and ZrO2 (3 mol% Y2O3) 1623 ceramics, which reveals that the appearance of c- phase results in the lowering of the transformation 4473K temperature 5423K 3.2 Kinetics of the isothermal/t-m transformation in ZrO2 3mo/%Y2O3) ceramics v8z Thermal expansion curves obtaincd at a hcating and cooling rate of 100C/min for ZrO2(3 mo 1%Y2O3)ceramics held at 200 C for 4 h are shown in Fig. 2:(a) refers to the curve of expansion vs temperature;()refers to the curve of expansion Time(s) time. Results of thermal expansion and XRD measurements for specimens held at difierent tem Fig. 3. Chan the fraction of m-phase with holding time for Zro2(3 mol%Y2O3) specimens held at different tempera peratures for 4 h are listed in Table 2, from which it Ires for 4 h

Isothermal tetragonal to monoclinic transformation 31 Table 1. Starting and ending temperatures of the t-+m and m+t transitions for Zr02(3 mol% Y20J) and Zr02(2 - mol%Y203) specimens sintered at 1600°C Specimens Zr0,(3mol%Y203) Zr02(2 mol%Y203) MS Mf AS Ar 387°C 142°C 374°C 542°C 459°C 322°C 567°C 602” C 0.2 - I I I I I I I 0 100 200 300 400 500 600 700 Temperature (“C) 0.8 -(b) 0.6 - 0.4 - 0.2 0 Time (ks) Fig. 2. Thermal expansion curves obtained at a heating and cooling rate lOO”C/min for ZrOz(3 mol% Y&I,) ceramics held at 200°C for 4 h: (a) curve of expansion vs temperature; (b) curve of expansion vs time. Table 2. Results of thermal expansion and XRD measurements for Zr02(3mol%Yz03) specimens held at differ￾ent temperatures for 4 h Temperature (“C) Pre-m phase Isothermal m-phase Athermal m-phase Overall m-phase Expansion before holding Expansion during holding Expansion after holding Overall expansion 150 15.53% 16.40% 31 .Q& 0.3532% 1.0460% 0 1.3990% 200 13.4% 52.47% 0 65.80% 0.2344% 0.9214% 0 1.1510% 250 10.93% 49.68% 2.69% 63.30% 0.1995% 0.9069% 0.04919% 1.1510% 300 13.23% 48.19% 3.20% 64.40% 0.24335% 0.8864% 0.0589% 1 .I 880% 350 3.60% 15.70% 7.20% 26.50% 0.2061% 0.4098% 0.4098% 1.5120% diagram. l2 The microstructure of the t + m dual phase obtained at room temperature implies that part of the sintered t-phase has transformed to m￾phase during cooling. However, ZrOz(3 mo- 1% Y203) ceramics were sintered in the c + t dual phase region and the microstructure obtained at room temperature is composed of c, t and m-pha￾ses. Table 1 lists the transformation temperatures of Zr02(2 mol% Y203) and Zr02(3 mol% Y203) ceramics, which reveals that the appearance of c￾phase results in the lowering of the transformation temperature. 3.2 Kinetics of the isothermal t-m transformation in Zr02(3 mol% YzOs) ceramics Thermal expansion curves obtained at a heating and cooling rate of lOO”C/min for Zr02(3mo- 1% Y203) ceramics held at 200°C for 4 h are shown in Fig. 2: (a) refers to the curve of expansion vs temperature; (b) refers to the curve of expansion vs time. Results of thermal expansion and XRD measurements for specimens held at different tem￾peratures for 4 h are listed in Table 2, from which it is evident that a cooling rate of lOO”C/min is not high enough to inhibit the t-m transformation before holding, and the lower the holding tem￾perature, the more pre-m phase produced prior to holding. Therefore, the initial phase constitutent for the measurement of the isothermal kinetics I 623K 2 573K 3 523K 4 473K 5 423K Time (s) Fig. 3. Changes of the fraction of m-phase with holding time for Zr02(3 mol% Y203) specimens held at different tempera￾tures for 4 h

w.Z. Zhu 673 Fig. 4. Time-temperature-transformation (TTT)curves for the(a)ZrO2(2 mol%Y2O3) and (b)ZrO2(3 mol%Y2O3)specimens curve consists of t and m phases, and part of the (c) The functional relationship between In(nk phase transforms isothermally to m-phase during and 1/T, shown in Fig. 6, reveals the activa- holding. Change of volume fraction of m-phase tion energy for the transformation to be obtained during holding with time is shown in 2274 kJ/ mol, irrespective of the holding Fig. 3, and the corresponding time-temperature- transformation (TTt) curve turns out to be C (d)The isothermal transformation cannot pro- shaped with a"nose"temperature determined to ceed to completion with part of the con be300° hown in Fig. 4. For comparison, the strained t-phase remaining in the final TTT curve of ZrO=(2 mol%Y2O3)ceramics is also microstructure shown in Fig. 4. The kinetics of the isothermal t-m transformation possesses the following fea As stated above, ZrO2 (3 mol%Y2O3)ceramics were fabricated in the c+t dual phase region at a sintering temperature of 1600oC, according to the (a)The existence of an incubation period which binary ZrO2-Y2O3 phase diagram '2 Morphologies initially becomes shorter with a decrease in of the c-phase grains are comparatively easier to holding temperature, then becomes longer distinguish from those of the t-phase grains. It was when the holding temperature is further previously reported that c-phase grains are rela decreased,resulting in the appearance of a tively difficult to nucleate during sintering and once nose"temperature. The existence of an they nucleate, growing speed is much higher than incubation period is one of the features of a that of t-phase grains until the phase equilibrium diffusional transformation whose isothermal state, which can be ascribed to the fact that a small kinetics can be expressed in terms of the amount of c-phase grains are primarily located at Johson-Mehl-Avrami (JMA)equation, 1.e f=1-exp(kr), where f is the volume frac- tion of transformed phase, k is a variable 1623K associated with the energy barrier for critical 2573K nucleation and growth, and n is a constant 4473K depending on the nucleation sites. The TTT 5423K curve of ZrO2 (3 mol%Y2O3)ceramics lies to 2 he down-left side of that of Zro2 (2 mo- 1%Y2O3)ceramics, indicating that at the same holding temperature, the incubation period of the former is shorter than that of he lat (b)The value of n, the exponent in the JMA equation, is somewhat different at different holding temperatures, implying slight varia tions of nucleating sites with holding tem- Fig. 5. Relationship between InIn 1/(1-) and Int for Zro2(3 perature(Fig. 5) mol% Y2O3)specimens

38 , W. Z. Zhu (a) 723- 523- 1% 10% I I I I 10 102 IO3 Time (s) 623- (b) 5 513 - 0 2 I & 523- F + 473- 423- I- IO 102 I03 IO“ Time (s) Fig. 4. Time-temperature-transformation (TTT) curves for the (a) ZrOz(2 mol% YzO3) and (b) ZrOT(3 mol% YzO3) specimens. curve consists of t and m phases, and part of the t￾phase transforms isothermally to m-phase during holding. Change of volume fraction of m-phase obtained during holding with time is shown in Fig. 3, and the corresponding time-temperature￾transformation (TTT) curve turns out to be C￾shaped with a “nose” temperature determined to be 300°C as shown in Fig. 4. For comparison, the TTT curve of Zr02(2 mol% Y203) ceramics is also shown in Fig. 4. The kinetics of the isothermal t+m transformation possesses the following fea￾tures: (a> co The existence of an incubation period which initially becomes shorter with a decrease in holding temperature, then becomes longer when the holding temperature is further decreased, resulting in the appearance of a “nose” temperature. The existence of an incubation period is one of the features of a diffusional transformation whose isothermal kinetics can be expressed in terms of the Johson-Mehl-Avrami (JMA) equation, i.e. f= 1 -exp(-kt”), where f is the volume frac￾tion of transformed phase, k is a variable associated with the energy barrier for critical nucleation and growth, and n is a constant depending on the nucleation sites. The TTT curve of Zr02(3 mol% Yz03) ceramics lies to the down-left side of that of Zr02(2mo- 1% Y203) ceramics, indicating that at the same holding temperature, the incubation period of the former is shorter than that of the latter; The value of n, the exponent in the JMA equation, is somewhat different at different holding temperatures, implying slight varia￾tions of nucleating sites with holding tem￾perature (Fig. 5). (c) The functional relationship between ln(nk) and l/T, shown in Fig. 6, reveals the activa￾tion energy for the transformation to be 22.74 kJ/mol, irrespective of the holding temperature. (d) The isothermal transformation cannot pro￾ceed to completion with part of the con￾strained t-phase remaining in the final microstructure. As stated above, Zr02(3 mol% Y203) ceramics were fabricated in the c + t d.ual phase region at a sintering temperature of 16OOC, according to the binary ZrO2-Y203 phase diagram.12 Morphologies of the c-phase grains are co’mparatively easier to distinguish from those of the t-phase grains. It was previously reported that c-phase grains are rela￾tively difficult to nucleate during sintering and once they nucleate, growing speed is much higher than that of t-phase grains until the phase equilibrium state, which can be ascribed to the fact that a small amount of c-phase grains are primarily located at 1 623K 2 573K -I- 3523K 4 473K Int Fig. 5. Relationship between lnln 1/(1-f) and In t for Zr02(3 - mol% Y204 specimens

Isothermal tetragonal to monoclinic transformation ceramics. This is because the "nose"'temperature of the TTt curve is largely determined by the starting temperature of the t-m transformation 3.3 Microstructure of ZrO2 3 mo/%Y203)ceramics Microstructure and morphologies of the m-phas are shown in Fig. 8, among which Fig 8(a) indi cates that the grain with dark contrast is monocli- nic phase, while the grain with relatively bright m dual phase. Figure 8(b)shows the different morphologies of m-phase within two irre l/Tx10+3 gular t-phase grains. In one grain the m-phase takes the form of parallel laths, which is similar to Fig. 6. Functional relationship between In(nk)and 1/T for the morphology of low-carbon martensite. Lath ZrO2(3 mol%Y2O3)specimens shaped m-phase is one of the features of the microstructure in ZrO2(Y2O3)ceramics, which is different from those in pure zirconia. In another the grain boundaries t-phase. 12 As a result, grain, the m-phase appears in the form of an"N the size of c-phase grai lways larger than that shape, large"N"typed m-phase passes across the of t-phase grains. SEM photographs of ZrO2 grain and the un-transformed part of the grain (2 mol%Y203) and Zro2 (3 mol% Y2O3)ceramics intercepted by a smaller one. " N"typed re simultaneously shown in Fig. 7, which indicates morphology is similar to that of high-carbon mar- that the average grain size of t-phase in the former tensite. The different morphology of the m-phase is Is larger than that in the latter with a small amount probably associated with different mechanisms of f larger c-phase grains. Therefore, formation, in that"N"typed m-phase is indicative preferential growth of c-phase grains inhibits the of auto-catalytic nucleation and coordinating growth of t-phase grains in ZrO2(3 mol%Y2O3) growth and the whole process is much faster, while ceramics and the existence of a small amount of c- the formation of lath-shaped m-phase is relatively phase grains refines the t-phase grains. The fact slower. Small triangular domains formed at inter chat the incubation periods of the ttt curves of sections between lattice invariant shear(Lis)twins ZrO2 (3 mol%Y2O3)ceramics are shorter than and deformation twins are shown in Fig 8(c).It those of the Ttt curves of Zro2(2 mol%Y203) significant that microcracks caused by collisions ceramics can be attributed to the refinement of the bctween the two kinds of twins appear at areas phase grains through the existence of c-phase without triangular domains and vice versa, indi grains. The lower"nose"temperature of the TTt cating that the triangular domain is essentially a curve of ZrO2 (3 mol% Y203)ceramics as com- type of coordinating twin, among which the LIS pared to that of zrO2(2 mol%Y2O3)ceramics can twin plane is(001)m and the twin plane of the be explained by the lower starting temperature of triangular domain is(011)m 13 Figure 8(d)illus- the t-m transformation of zro2(2 mol% Y203) trates the case in which the m- plate is deflected by a Fig. 7. SEM photographs showing the natural surfaces of (a)ZrO2(2 mol%Y203)and(b)ZrO2(3 mol%Y203)specimens

Isothermal tetragonal to monoclinic transformation l/T x 1O‘4 (K-l) Fig. 6. Functional relationship between In(&) and l/T for ZrOz(3 mol% YzO3) specimens. the grain boundaries of the t-phase.12 As a result, the size of c-phase grains is always larger than that of t-phase grains. SEM photographs of Zr02 (2 mol% Y203) and ZrOz(3 mol% Y203) ceramics are simultaneously shown in Fig. 7, which indicates that the average grain size of t-phase in the former is larger than that in the latter with a small amount of larger c-phase grains. Therefore, in some sense, preferential growth of c-phase grains inhibits the growth of t-phase grains in Zr02(3mol% Y203) ceramics and the existence of a small amount of c￾phase grains refines the t-phase grains. The fact that the incubation periods of the TTT curves of Zr02(3 mol% YzOs) ceramics are shorter than those of the TTT curves of Zr02(2mol% Y203) ceramics can be attributed to the refinement of the t-phase grains through the existence of c-phase grains. The lower “nose” temperature of the TTT curve of Zr02(3mol% Y203) ceramics as com￾pared to that of Zr02(2 mol% Y203) ceramics can be explained by the lower starting temperature of the t--+m transformation of Zr02(2 mol% Y203) 39 ceramics. This is because the “nose” temperature of the TTT curve is largely determined by the starting temperature of the t-+m transformation. 3.3 Microstructure of ZrO2(3mo/% Y2Os) ceramics Microstructure and morphologies of the m-phase are shown in Fig. 8, among which Fig. 8(a) indi￾cates that the grain with dark contrast is monocli￾nic phase, while the grain with relatively bright contrast is t + m dual phase. Figure 8(b) shows the different morphologies of m-phase within two irre￾gular t-phase grains. In one grain the m-phase takes the form of parallel laths, which is similar to the morphology of low-carbon martensite. Lath￾shaped m-phase is one of the features of the microstructure in ZrOz(Y20s) ceramics, which is different from those in pure zirconia. In another grain, the m-phase appears in the form of an “N” shape, large “N” typed m-phase passes across the grain and the untransformed part of the grain is intercepted by a smaller one. “N” typed m-phase morphology is similar to that of high-carbon mar￾tensite. The different morphology of the m-phase is probably associated with different mechanisms of formation, in that “N” typed m-phase is indicative of auto-catalytic nucleation and coordinating growth and the whole process is much faster, while the formation of lath-shaped m-phase is relatively slower. Small triangular domains formed at inter￾sections between lattice invariant shear (LIS) twins and deformation twins are shown in Fig. 8(c). It is significant that microcracks caused by collisions between the two kinds of twins appear at areas without triangular domains and vice versa, indi￾cating that the triangular domain is essentially a type of coordinating twin, among which the LIS twin plane is (OOl), and the twin plane of the triangular domain is (01 1),.13 Figure 8(d) illus￾trates the case in which the m-plate is deflected by a Fig. 7. SEM photographs showing the natural surfaces of (a) Zr02(2mol% Y203) and (b) ZrOz(3 mol% Y203) specimens

w.Z. Zhu dislocation during propagation, implying that the ture between c and t phases, and the larger the stress field of dislocation counteracts that of trans- effective driving force for precipitation;(2)nuclea formation tion of t-phase is kinetically favoured at regio Diffusional phase scparation to producc precipi- with relatively low yttria content because the con tates of t-phase has obviously occurred within centration of yttria is reduced in forming nuclei some regions of the c-phase grains during sintering, Comprehensively speaking, the phase constitu as shown ig. (a)and (b), respectively. tent in ZrO (3 mol%Y,O3)ceramics the Morphologies of the m-phase are quite different, in complicated and is composed of cubic, tetragonal that"N" typed m-phase exists in Fig 9(a) with and monoclinic phases. The t-phase can be either microcracks at the grain boundary, while m-phase sintered phase or precipitate depending on sinter- appears in the form of parallel laths in Fig 9(b) ing temperature and cooling conditions. The grain without microcracks at the grain boundary. a dif- size of the sintered t-phase is relatively small, while fraction pattern along the [lll] direction is shown precipitates of t-phase are produced in the form of in Fig. 9(c), in which three(112)type forbidden a"colony" microstructure through diffusional spots appear. Figure 9(d) shows the dark field phase separation within large c-phase grains during image taken by using a(112)forbidden spot, in high temperature sintering. The m-phase can either colony "microstructure- which is a occupy the whole grain or co-exist with the t-phase characteristic of product of the c-t diffusional within the original t-phase grains phase separation-is clearly observed Every(1 12 reflection corresponds to one kind of"colony" 3. 4 Mechanism of the isothermal t-m transition variant, each of which is composed of two groups of twinned precipitating t-phase with planes deter In tetragonal ZrOz, zirconium ions occupy sites of mined to be(101)m. It is found that regions with a the face-centred tetragonal lattice, where the dis lower yttria content in the c-phase grain are more tribution of oxygen ions deviates slightly from the favourable for the precipitation of t-phase, which (001)plane, leading to the appearance of certain can be accounted for as follows:()the lower the tetragonality. When Y2O3 is solid-solutioned into yttria content, the higher the equilibrium tempera- the Zro2 lattice, some of the zirconium ions are 0.2m 0.2um 0.2pm Fig.8. TEM photographs showing the microstructure of Zro2 (3 mol%Y203)specimens:(a)butterfly-like m-phase within a t phase grain and distribution of yttria content within different grains;(b) morphologies of m-phase within different grains; (c)tri- angular domains in twinned m-phase (d )interaction between the dislocations and the m-phase

40 dislocation during propagation, implying that the stress field of dislocation counteracts that of trans￾formation. Diffusional phase separation to produce precipi￾tates of t-phase has obviously occurred within some regions of the c-phase grains during sintering, as shown in Fig. 9(a) and (b), respectively. Morphologies of the m-phase are quite different, in that “N” typed m-phase exists in Fig. 9(a) with microcracks at the grain boundary, while m-phase appears in the form of parallel laths in Fig. 9(b) without microcracks at the grain boundary. A dif￾fraction pattern along the [l 1 l] direction is shown in Fig. 9(c), in which three (112) type forbidden spots appear. Figure 9(d) shows the dark field image taken by using a (112) forbidden spot, in which a “colony” microstructure - which is a characteristic of product of the c-+t diffusional phase separation - is clearly observed. Every (112) reflection corresponds to one kind of “colony” variant, each of which is composed of two groups of twinned precipitating t-phase with planes deter￾mined to be { lOl},. It is found that regions with a lower yttria content in the c-phase grain are more favourable for the precipitation of t-phase, which can be accounted for as follows: (1) the lower the yttria content, the higher the equilibrium tempera￾W. Z. Zhu ture between c and t phases, and the larger the effective driving force for precipitation; (2) nuclea￾tion of t-phase is kinetical1.y favoured at regions with relatively low yttria content because the con￾centration of yttria is reduce’d in forming nuclei. Comprehensively speaking, the phase constitu￾tent in Zr02(3 mol% Y203) ceramics is rather complicated and is composed of cubic, tetragonal and monoclinic phases. The t-phase can be either sintered phase or precipitate depending on sinter￾ing temperature and cooling conditions. The grain size of the sintered t-phase is relatively small, while precipitates of t-phase are produced in the form of a “colony” microstructure through diffusional phase separation within large c-phase grains during high temperature sintering. The m-phase can either occupy the whole grain or co-exist with the t-phase within the original t-phase grains. 3.4 Mechanism of the isothermal t-m transition In tetragonal Zr02, zirconium ions occupy sites of the face-centred tetragonal lattice, where the dis￾tribution of oxygen ions deviates slightly from the (001) plane, leading to the appearance of certain tetragonality. When Y203 is solid-solutioned into the ZrO* lattice, some of the zirconium ions are Fig. phar 8. TEM photographs showing the microstructure of Zr02(3mol% YzOs) specimens: (a) butterfly-like m-phase within a t- ;e grain and distribution of yttria content within different grains; (b) morphologies of m-phase within different grains; (c) tri￾angular domains in twinned m-phase; (d) interaction between the dislocations and the m-phase

Isothermal tetragonal to monoclinic transformation substituted for yttrium ions, during which a certain where Vo is the oxygen vacancy. The amount of re ected to be oxygen vacancies can be characterized by following produced to hold the ionic neutralization. The fol- formula: 4 wing reaction can be used to describe above process: 8 ZrO2+Y203-2Yzr vo+ 300 100+M (d) 0.3 Fig9. Precipitation of t-phase within a c-phase grain of Zro2(3 mol%Y203)specimens:(a)and(b) bright field images showing ne co-existence of c, t and m-phases;(c)diffraction pattern of the c-phase; (d)dark field image taken by using a(112)forbidden 0.3pm 0.3m Fig. 10. TEM photographs showing the in-situ formation of parallel lathed m-phase, induced by the irradiation of an electron beam in a ZrO, (2 mol%Y2O3) ceramic

Isotherm !a1 tetragonal to monoclinic transformation substitu ted for yttrium ions, during which a certain number of oxygen vacancies are expected to be produce :d to hold the ionic neutralization. The fol￾lowing reaction can be used to describe above process: 8 Zr02 + Y2O3 + 2Yz, + Vo + 300 41 where V0 is the oxygen vacancy. The amount of oxygen vacancies can be characterized by following formula: l4 v= 4M lOO+h4 I, 0.3 pm I Fig. 9. Precipitation of t-phase within a c-phase grain of Zr02(3 mol% YzOs) specimens: (a) and (b) bright field images showing the co-existence of c, t and m-phases; (c) diffraction pattern of the c-phase; (d) dark field image taken by using a (112) forbidden spot. Fig. 10. TEM photographs showing the in-situ formation of parallel lathed m-phase, induced by the irradiation of an electron beam in a ZrOz(2 mol% Y*Os) ceramic

w.Z. Zh Here y is the volume fraction of oxygen vacancies 4 CONCLUSIONS and M is the concentration of Y, O in mol%. It can be seen that the concentration of oxygen (a) The time-temperature-transformation TTT) vacancies incrcascs with increasing Y2 O3 concen curve for the t→→ m transformation in Thus, it is proposed that in ZrO2(Y2O3) ZrO2(3 mol% Y2O3)ceramics is C-shaped, containing a large amount of vacancies, the trans- with a"nose"temperature determined to be formation behaviour is quite different from that in 300C. The TTt curve of Zro2 (3mo- pure zirconia > For pure zirconia, when the t-m Io Y2O3)ceramics lies to the down-left side transformation occurs, the lattice composed of zir of that of Zro2 (2 mol%Y2O3)ceramics conium ions shears in a coordinating and military (b) Preferential growth of manner, concurrent with obvious volume expan- ZrO2 (3 mol%Y203) ceramics effectively sion as well as shape change. As a result, this dif- refines the grain size of the t-phase fusionless shear process becomes a predominant (c) The microstructure of Zro2 (3 mol%Y2O3) factor in controlling the kinetics of transformation, ceramics is very complicated, in that it con- which appears to be athermal. Because ceramic tains tetragonal phase (which can either be a materials possess high strength and elastic modu sintered phase or precipitates), monoclinic lus, the strain energy caused by a shear-like trans phase(which can either occupy the whole formation is large and morphologies of the final grain or co-exist with the tetragonal phase products, which are always twinned, are deter and cubic phase mined by the strain energy For ZrO2(Y203)cera (d)The mechanism of the t-m transition in mics. when the t-m transformation occurs, not Zro2(Y2O3)ceramics is different from that in only a coordinating shift of zirconium ions is pure ZrO2. The nucleation and longitudinal ded but short-range diffusion of oxygen ions also growth of m-phase plates in ZrO2(Y2O3 takes place, with the implication that this kind of transformation possesses both displacive and non growth thereafter is controlled by short displacive features. TEM photographs of the in- range diffusion of oxygen ions and, in this situ growth of m-phase induced by the irradiation the t-m transformation of an electron heam are shown in Fig. 10, from ZrO2(Y2O3)ceramics possesses both displa which direct proof of heterogeneous nucleation of cive and non-displacive features m-plates at the grain boundaries can be obtained It is observed that the nucleation process is rather rapid and once nucleation completes, growth in the REFERENCES longitudinal direction is much faster than that in the tranverse direction. It is roughly estimated th CHEN, I w.& CHIAO, Y H, Acta Metall, 33(10) the growing velocity along the longitudinal direc- 2 HEuEr. AH. clausSEN N. KRIVEN w.M.& tion is 30 times faster than that along the tranverse RUHLE. M.J. Am Ceram Soc., 65(12)(1982)642 direction. It is thus proposed that the nucleation 3. SUBBARAO, E. C, Advances in and longitudinal growth of m-phase plates in Ceramics Society, Columbus, OH, 1981, p. I ZrO2(Y2O3)ceramics is believed to be displacive, 4. SATO, T, OHTEKL, S, ENDO, T& SHIMADA, M while the sidewise growth thereafter is proved to be short-range diffusion controlled. The short-range 5. NAKANISHI, N &SHIGEMATSU,T,Mater.Trans diffusion of oxygen ions becomes a predominant 6. LEL, T. C, ZhU, w.z.& ZhoU, Y, Mater. Chem factor in controlling the kinetics of the t-+m Phys.34(4)(1993)317 transformation in ZrO2(Y2O3)ceramics. The 7. ZHU, W.Z., LEl, T C.& ZHOU, Y, J. Mater. Sci release of the large transformation strain incurred 8(12)(1993)6479 8. NAKANISHi, N. SHIGEMATSU, T Muter. Truns. by the expansion of monoclinic phase through the JM,32(8)(1991)778. short-range diffusion of oxygen ions might be 9. nakanishi n.& shigematsu. t. zirconia responsible for the diversity of m-phase morphol Ceram,8(1)(1986)71 10. BEHRENS. D. DRANSMANNG w& heuer. a ogies observed in ZrO2(Y2O3)ceramics. In view of J.Am. ceran.Soc,76(4)(1993)1025 the fact that the transformation activation energ 11. GARVIE, R. C, J. Am. Ceram. Soc., 55(6)(1972) calculated using the kinetics data is in the range of Advances in Ceramics, Science 20-40 kJ/mol, which is far less than the activation IL, Vol. 12. American Ceramics Society, C energy for self-diffusion of oxygen ions (96 kJj OH,1984,pp.352 mol), it is speculated that the shifting distance of 13. WOLTEN, G. M, J. Am. Ceram. Soc., 46(10)(1963) the oxygen ions during transition is less than the 14. INGEL R. P.& Ill. D. L.J. Am. Ceram. Soc., 69(4) lattice constant of the tetragonal phase(0.50 nm) (1986)325

42 Here V is the volume fraction of oxygen vacancies and M is the concentration of Y203 in mol%. It can be seen that the concentration of oxygen vacancies increases with increasing Y203 concen￾tration. Thus, it is proposed that in Zr02(Y203) containing a large amount of vacancies, the trans￾formation behaviour is quite different from that in pure zirconia. is For pure zirconia, when the t-+m transformation occurs, the lattice composed of zir￾conium ions shears in a coordinating and military manner, concurrent with obvious volume expan￾sion as well as shape change. As a result, this dif￾fusionless shear process becomes a predominant factor in controlling the kinetics of transformation, which appears to be athermal. Because ceramic materials possess high strength and elastic modu￾lus, the strain energy caused by a shear-like trans￾formation is large and morphololgies of the final products, which are always twinned, are deter￾mined by the strain energy. For Zr02(Y203) cera￾mics, when the t--+m transformation occurs, not only a coordinating shift of zirconium ions is nee￾ded, but short-range diffusion of oxygen ions also takes place, with the implication that this kind of transformation possesses both displacive and non￾displacive features. TEM photographs of the in￾situ growth of m-phase induced by the irradiation of an electron beam are shown in Fig. 10, from which direct proof of heterogeneous nucleation of m-plates at the grain boundaries can be obtained. It is observed that the nucleation process is rather rapid and once nucleation completes, growth in the longitudinal direction is much faster than that in the tranverse direction. It is roughly estimated that the growing velocity along the longitudinal direc￾tion is 30 times faster than that along the tranverse direction. It is thus proposed that the nucleation and longitudinal growth of m-phase plates in Zr02(Y203) ceramics is believed to be displacive, while the sidewise growth thereafter is proved to be short-range diffusion controlled. The short-range diffusion of oxygen ions becomes a predominant factor in controlling the kinetics of the t+m transformation in ZrOz(YzO3) ceramics. The release of the large transformation strain incurred by the expansion of monoclinic phase through the short-range diffusion of oxygen ions might be responsible for the diversity of m-phase morphol￾ogies observed in ZrOz(Y203) ceramics. In view of the fact that the transformation activation energy calculated using the kinetics data is in the range of 2640 kJ/mol, which is far less than the activation energy for self-diffusion of oxygen ions (96 kJ/ mol), l6 it is speculated that the shifting distance of the oxygen ions during transition is less than the lattice constant of the tetragonal phase (0.50 nm). W. Z. Zhu 4 CONCLUSIONS (4 cc> (4 The time-temperature--transformation (TTT) curve for the t-+m transformation in Zr02(3 mol% Y203) c,eramics is C-shaped, with a “nose” temperature determined to be 300°C. The TTT curve of Zr02(3mo- 1% Y203) ceramics lies to the down-left side of that of Zr02(2 mol% Y20s)ceramics. Preferential growth of c-phase grains in Zr02(3 mol% Y20~) ceramics effectively refines the grain size of the t-phase. The microstructure of Zr02(3 mol% Y203) ceramics is very complicated, in that it con￾tains tetragonal phase (which can either be a sintered phase or precipitates), monoclinic phase (which can either occupy the whole grain or co-exist with the tetragonal phase) and cubic phase. The mechanism of the t+m transition in Zr02(Y203) ceramics is different from that in pure Zr02. The nucleation and longitudinal growth of m-phase plates in Zr02(Y203) ceramics are displacive, while the sidewise growth thereafter is controlled by short￾range diffusion of oxygen ions and, in this sense, the t--tm transformation in ZrOz(Yz03) ceramics possesses both displa￾cive and non-displacive features. REFERENCES 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. CHEN, I. W. & CHIAO, Y H., Acta Mefall., 33(10) (1985) 1827. HEUER, A. H., CLAUSSEN, N., KRIVEN, W. M. & RUHLE, M., J. Am. Ceram. Sot., 65(12) (1982) 642. SUBBARAO, E. C., Advance.x in Ceramics, Science and Technology of Zircon& Vol. 3. American Ceramics Society, Columbus, OH, 1981, p. 1. SATO, T., OHTEKI, S., ENDO, T. & SHIMADA, M., J. Am. Gram. SOL, 68(10) (1985) C-320. NAKANISHI, N. & SHIGEMATSU, T., Mater. Trans. JIM, 33(3) (1992) 318. LEI, T. C., ZHU, W. Z. & ZHOU, Y., Mater. Chem. Phys., 3q4) (1993) 317. ZHU, W. Z., LEI, T. C. & ZHOU, Y., J. Mater. Sci., 28(12) (1993) 6479. NAKANISHI, N. & SHIGEMATSU, T., Mater. Trans. JIM, 32(8) (1991) 778. NAKANISHI, N. & SHIGEMATSU, T., Zirconiu Gram., 8(l) (1986) 71. BEHRENS, D., DRANSMANN, G. W. & HEUER, A. H., J. Am. Ceram. Sot., 76(4) (1993) 1025. GARVIE, R. C., J. Am. Ceram. Sot., 55(6) (1972) 303. RUHLE, M., CLAUSSEN, N. & HEUER, A. H., Advances in Ceramics, Science and Technology of Zirco￾niu ZI, Vol. 12. American Ceramics Society, Columbus, OH, 1984, pp. 352. WOLTEN, G. M., J. Am. Ceram. Sot., 46( 10) (1963) 418. INGEL, R. P. & III, D. L., .I. Am. Ceram. Sot., 69(4) (1986) 325

Isothermal tetragonal to monoclinic transformation 15. BANSAL, G K& HEUER, A H, Acta Metall, 20(11) 16. TSUBAKINO, H, SODONA, K& NOZATO,R,J (1972)1281. Mater. Sci. Le,12(3)(1993)196

Isothermal tetragonal to monoclinic transformation 43 15. BANSAL, G. K. & HEUER, A. H., Acta Metall., 20(11) 16. TSUBAKINO, H., SODONA, K. & NOZATO, R., J. (1972) 1281. Mater. Sci. Lett., 12(3) (1993) 196