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Effect of minor SiO2 addition on the creep behavior of superplastic tetragonal ZrO2
Acta Materialia 52 (2004) 3355–3364 www.actamat-journals.com
Effect of minor SiO2 addition on the creep behavior of superplastic tetragonal ZrO2 K. Morita *, K. Hiraga, B.-N. Kim National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan Received 2 December 2003; received in revised form 22 March 2004; accepted 24 March 2004 Available online 21 April 2004
Abstract A sigmoidal stress–creep rate relationship appearing in tetragonal ZrO2 drastically changes with SiO2 addition, where silicon segregates along the grain boundaries and dissolves into the ZrO2 grains. This change is completed with the precipitation of glass pockets though the creep rate gradually increases with increasing the volume fraction. The glass pockets merely increase the creep rate by accelerating grain boundary sliding (GBS) by the viscous flow of the glass phase. Although the silicon segregated along the grain boundaries may enhance GBS, the solute silicon seems to control the rate of creep. The solute silicon increases the rate of Nabarro–Herring creep by accelerating the lattice diffusivity of cations at low stresses but decreases the rate of GBS by inhibiting the accommodation process at high stresses. Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Ceramics; High-temperature deformation; Superplasticity; Tetragonal ZrO2
1. Introduction Silicate glasses have deliberately been added to ZrO2 ceramics because they can enhance sintering [1,2]. The addition of glasses is known to play an important role in the mechanical properties of ZrO2 ceramics, particularly at elevated temperatures [3–7]. The addition of pure SiO2 can enhance the superplasticity of fine-grained Y2 O3 -stabilized tetragonal ZrO2 polycrystal (Y-TZP) [6,7]; e.g., 5 wt%-SiO2 -doped Y-TZP exhibits a large tensile ductility exceeding 1000% at 1673 K [6]. The enhanced superplasticity of SiO2 doped Y-TZP has been attributed to plastic flow accelerated through glass pockets [6] and a change in the chemical bonding state of grain boundaries caused by the segregation of silicon [8]. The former factor can enhance the accommodation process of stress concentrations exerted at multiple grain junctions by grain boundary sliding (GBS) and the latter factor can strengthen the bonding of the grain boundaries. *
Corresponding author. Tel.: +81-29-859-2537; fax: +81-29-8592501. E-mail address: [email protected] (K. Morita).
Earlier studies [9–15] have noted that the superplastic flow behavior of Y-TZP is sensitive to the existence of residual impurities. This is because the impurities significantly affect GBS that has generally been accepted as the primary flow mechanism. For a large amount of SiO2 addition, both factors simultaneously work on the superplastic flow and hence it is hard to examine the respective roles of SiO2 addition. For understanding the exact role of SiO2 addition in the superplasticity of Y-TZP, the deformation behavior should be examined over a wide range of SiO2 addition with and without glass phase precipitation. The present study was therefore performed to examine the effect of SiO2 on the creep behavior of finegrained Y-TZP as a function of the SiO2 content. The creep behavior was correlated with the microstructure varied with the SiO2 content ranging from 0 to 2.5 wt%.
2. Experimental procedures The materials used were prepared from a high-purity 3Y-TZP powder containing 3 mol% Y2 O3 (Tosoh Co., Ltd.,: 50 SiO2 , <50 Al2 O3 , <50 Fe2 O3 , 220 NaO in wt
1359-6454/$30.00 Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2004.03.033
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ppm) and pure SiO2 (Nissan Chemical Industry Co., Ltd.) using a method employed by Kajihara et al. [5]. The 3Y-TZP powders doped with 0–2.5 wt% SiO2 were cold isostatically pressed at about 200 MPa, sintered at 1673 K for 3 h in air and cooled in the furnace. The sintered materials had a relative density higher than 98%. The initial grain size, d0 , defined as 1.56 times the average intercept lengths of the grains [16] was about 0.35 lm. From the sintered bodies, dog-bone-shaped flat tensile specimens were machined with a gauge portion of t w l 3- 3- 20 mm. Tensile creep tests were conducted at 1673 K in air and at initial stresses of 0.8–80 MPa. We corrected the interference arising from the deformation around the grip sections and grain growth using a method described elsewhere [17,18]. Briefly, true displacement was directly measured by monitoring the length between the targets made at both ends of the gauge portion. The monitored creep rates, e_ m , can be corrected for instantaneous stress as e_ r ¼ e_ m expðnea Þ;
ð1Þ
where ea is the arbitrary strain. When grain growth occurs during deformation, Eq. (1) can be corrected for the instantaneous grain size, da , at ea as e_ rd ¼ e_ m expðnea Þðd0 =da Þp :
ð2Þ
The as-sintered and deformed microstructures were examined by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Before the SEM observations, the specimens were mechanically polished and thermally etched at 1473 K for 30 min. For the TEM observation of the as-sintered microstructures, several specimens were quenched from 1673 K by di-
rectly dropping from the furnace into water so that they must preserve the microstructure at high-temperature. Sheets with a thickness of 500 lm were cut with a lowspeed diamond cutter, mechanically polished to less than 100 lm in thickness and further thinned with an Ar ion-milling machine. The energy dispersive X-ray spectroscopy (EDS) was performed at areas thin enough to obtain lattice image using a JEOL-2010F with a probe size of about 1.0 nm.
3. Experimental results 3.1. As-sintered microstructure As shown in Fig. 1, the microstructure of the as-sintered materials depends on the SiO2 content and can be classified into two groups. For a small amount of the SiO2 addition less than 0.3 wt%, no glass phase is found at any grain junction, but beyond this, nanometer-sized glass pockets are apparently observed to precipitate at rounded multiple grain junctions as indicated by the triangles. At P 0.3 wt% SiO2 , the volume fraction of the glass pockets monotonously increases with the SiO2 content. Further details of the microstructure were examined using high-resolution TEM and EDS in Fig. 2. EDS analysis clearly revealed that the nanometer-sized glass pockets in the 0.3–2.5 wt%-SiO2 -doped materials consisted mainly of SiO2 . For the undoped materials, the segregation of yttrium was found along the grain boundaries, whereas for the SiO2 -doped materials, silicon was found to segregate along the grain boundaries
Fig. 1. Bright-field TEM images of the as-sintered 3Y-TZP doped with (a) 0, (b) 0.07, (c) 0.3 and (d) 2.5 wt% SiO2 .
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On the other hand, EDS spectra obtained from the grain interior revealed a trace of silicon in addition to zirconium and yttrium (Fig. 2(c)). Since the EDS measurements were carried out at sufficient distances, at least >20 nm, from the grain boundaries and glass pockets, the interference from those areas can be negligible even if beam scattering occurred through thin foils. This shows that a small amount of silicon also dissolves into ZrO2 grains. The level of silicon along the grain boundaries and within the grains was deduced from the X-ray intensity of the Si–K peak in the EDS spectra. As shown in Fig. 3, the level of silicon depends on the SiO2 content. For <0.3 wt% SiO2 , the level steeply increases with the SiO2 addition, whereas for P 0.3 wt% SiO2 , it is almost constant independently of the SiO2 content. The saturation value is compatible with earlier data for 5 wt%-SiO2 -doped 2.5Y-TZP [8]. Although the level of silicon along the grain boundaries is apparently higher than that of the grain interiors, it tends to be lower in the water-quenched materials than in the furnace-cooled ones. 3.2. Tensile creep behavior The creep rate curve of the present materials strongly depends on the correction of grain size as shown in Fig. 4. For the e_ m and e_ r curves, the creep rates gradually change with e and no well-defined steady-state region appears. For the e_ rd curves, on the other hand, a well-defined steady-state region appears at e P 0:03. It is noticeable that the discrepancy between the e_ r and e_ rd curves increases with e and it becomes more prominent at lower stresses; it exceeds a factor of 2 at e P 0:1 for 3 MPa. Thus, the occurrence of grain growth strongly affects the creep behavior of fine-grained 3Y-TZP. Since 0.20
Fig. 2. (a) High-resolution TEM image of a multiple grain junction filled with the glass phase and typical EDS spectra obtained from (b) grain boundaries and (c) grain interiors for the water-quenched 3YTZP doped with 0.3 wt% SiO2 .
in addition to yttrium as shown in Fig. 2(b). In the present materials, however, the lattice fringes of adjacent grains directly intersect at the boundaries and no amorphous phase was found along the grain boundaries as shown in Fig. 2(a). If an amorphous phase exists along grain boundaries, the dihedral angle at glass pockets must be less than 10° [19,20]. The large dihedral angles of 60° strongly support the fact that the grain boundaries of the SiO2 -doped materials are not wetted by amorphous phases even at high-temperatures [8].
Intensity Ratio, Si/Zr
furnace- watercooled quenched
0.15
Grain Boundary Grain Interior
Ikuhara et al. [8]
0.10
0.05
0
0
0.5
1.0
1.5
2.0
2.5
3.0
5.0
SiO2 Content, wt% Fig. 3. The X-ray intensity of the Si–K peak normalized by the Zr–L peak, Si/Zr, plotted as a function of the SiO2 content. The circles and squares represent the data of the furnace-cooled and water-quenched specimens, respectively. For comparison, earlier data for 5 wt%-SiO2 doped 2.5Y-TZP [8] are also plotted.
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(a) 3Y-TZP + 0.3wt%SiO2
the creep behavior should be characterized in the steadystate region, the grain size correction is necessary for the SiO2 -doped 3Y-TZP, as noted in the recent study for high-purity 3Y-TZP [17].
1673 K - 3 MPa
Creep Rate, ε / s-1
10-5
•
•
εσd •
ε•m εσ
10-6
(b) 3Y-TZP + 0.3wt%SiO2
1673 K - 30 MPa
-3
•
10
Creep Rate, ε / s-1
3.3. Stress–strain rate relationship
εm •
εσd
•
•
εσ
10-4 0
0.05
0.1
True Strain, ε
0.15
0.2
Fig. 4. Creep rate curves plotted as a function of true strain, e, for 0.3 wt%-SiO2 -doped 3Y-TZP obtained at 1673 K and at (a) 3 and (b) 30 MPa.
The steady-state creep rates, e_ rd , for 3Y-TZP without and with glass pockets are plotted as a function of true stress, r, in Fig. 5(a) and (b), respectively. The e_ rd –r relationship shows three prominent features. First, the e_ rd –r relationship is divided into three categories; a high-stress region (Region H) with n 1:8–2:7, an intermediate-stress region (Region I) with n 2:5–5:0 and a low-stress region (Region L) with n 1:3. Thus, the e_ rd –r relationship as a whole exhibits a sigmoidal feature in all the SiO2 contents examined though the sigmoidal feature becomes less noticeable for P 0.3 wt% SiO2 . Second, the sigmoidal e_ rd –r relationship is not a simple function of the SiO2 content. For <0.3 wt% SiO2 , the sigmoidal e_ rd –r relationship sensitively changes with the SiO2 addition (Fig. 5(a)), whereas for P 0.3 wt% SiO2 , the general feature is insensitive to the SiO2 addition though the creep rates gradually increase with the SiO2 content (Fig. 5(b)). Third, the effect of the SiO2 addition is different among Regions H, I and L. The n value is dependent on
10-1
Corrected Creep Rate, εσd / s-1
(a) without glass pocket
•
SiO2 (wt%) 0 [17]
10-2 10
(b) With glass pocket Kajihara et al. [6] 2.5Y-TZP+5wt%SiO2 D0 = 0.25µm
2.7
0.02 0.07
-3
1.8
10-4
1.8
2.5
10-5 1.3
2.5
5.0
0.07 0.3 0.9 2.5
10-6 1.3
10-7 Region L
10
Region I
Region H
Region L
SiO2 (wt%)
Region I
Region H
-8
1
10 100 True Stress, σ / MPa
1
10 100 True Stress, σ / MPa
Fig. 5. Corrected creep rates, e_ rd , plotted as a function of true stress, r, for 3Y-TZP doped with (a) 0–0.07 and (b) 0.07–2.5 wt% SiO2 . For comparison, earlier data for 5 wt%-SiO2 -doped 2.5Y-TZP [6] are also plotted. The open and closed symbols represent the creep data for the materials without and with glass pockets, respectively.
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the SiO2 content in Regions H and I, but independent of the content in Region L. In Regions H, for example, the SiO2 addition tends to decrease the creep rate and steeply decreases the n value from 2.7 to 1.8 up to 0.3 wt%, but beyond this, it slightly enhances the creep rate with maintaining n 1:8. In Region L, on the other hand, the SiO2 addition enhances the creep rate but does not change the n value of 1.3. According to Gust et al. [12], the addition of a small amount of barium silicate glass appears to increase the shear modulus, G, of 3Y-TZP. If this is also the case for the SiO2 -doped 3Y-TZP, the SiO2 -dependent creep behavior shown in Fig. 5 may be related to a change in the G value due to glass addition. However, the change in the G value is 4% [12] and negligible. In addition, the e_ rd –r relationship is not a simple function of the SiO2 content. Thus, the general feature of the SiO2 -dependent e_ rd –r relationship cannot be ascribed to a change in the G value accompanying the SiO2 addition. 3.4. Deformed microstructure Fig. 6 shows an example of a deformed microstructure. After deformation, the specimen was quickly cooled under loading to preserve the microstructures developed during deformation. Although the deformed materials almost retained the equiaxed grain shapes, strain contrasts appeared to be more pronounced after deformation. The co-segregation of silicon and yttrium was also detected along grain boundaries, but no intergranular amorphous phase was found even after deformation, as shown in Fig. 6(b). 1 Dark-field TEM observation reveals densely aligned intragranular dislocations, which seem to nucleate around multiple grain junctions indicated by an arrow in Fig. 6(c). Since these dislocation substructures were not observed before deformation, the dislocations are developed during deformation. 4. Discussion 4.1. Microstructural change due to SiO2 addition In earlier studies on glass-doped Y-TZP [5,8–14], it has been believed that the solubility of silicon into ZrO2 grains was highly limited, and hence, it formed an intergranular amorphous film with a thickness of about 1.0 nm. In contrast, the present study shows that a small amount of silicon dissolves into the ZrO2 grains. In 1
Glass phases precipitated at multiple grain junctions were observed to penetrate into two grain junctions only around the ahead of cracklike cavities, which are formed at e > 0:8 [20]. However, such crack-like cavities were not observed at e < 0:15, where the e_ rd values were determined, and no grain boundary amorphous phase was found in the present study.
Fig. 6. (a) Bright-field, (b) high-resolution and (c) dark-field TEM images for 2.5 wt%-SiO2 -doped 3Y-TZP after deformation up to e 0:18 at 1673 K and at 30 MPa.
addition, although the grain boundaries of the SiO2 doped materials are enriched with silicon, no amorphous phase occurs along the grain boundaries even after deformation. The present results are consistent with the earlier studies by Ikuhara et al. [8,21] and Aoki et al. [22], showing the segregation of silicon but no
K. Morita et al. / Acta Materialia 52 (2004) 3355–3364
amorphous phase along grain boundaries and the dissolution of silicon into ZrO2 grains. The level of silicon segregation along the grain boundaries tends to be higher in the furnace-cooled materials than in the water-quenched ones. According to the TEM in-situ heating experiment by Ikuhara et al. [21], at high temperatures, a small amount of SiO2 is soluble into the ZrO2 grains in the vicinity of the grain boundaries, but during cooling, the dissolved SiO2 precipitates from the super-saturated SiO2 –ZrO2 solid solution. This result suggests that the higher silicon segregation in the furnace-cooled materials may be caused by the precipitation of SiO2 , which dissolved into the ZrO2 grains at high temperatures. The recent study of Gremillard et al. [23] has noted that no silicon enrichment was observed along grain boundaries and no detectable silicon was found in grain interiors in 2.5 wt%-SiO2 -doped 3Y-TZP. Although the reason for the discrepancy is not clear at the present time, the conclusion clearly contradicts to the present microstructural examination, showing that no glass pocket was found for SiO2 additions less than 0.3 wt%. If silicon is insoluble into the ZrO2 matrix, glass pockets should precipitate simultaneously with the SiO2 addition. In addition, we confirmed that a decrease in the grain boundary area due to grain growth enhanced the precipitation of the glass pockets [20]. As a result, the change in the microstructure accompanying the SiO2 addition can be explained as follows. For a small amount of the SiO2 addition less than 0.3 wt%, the doped SiO2 preferentially segregates along the grain boundaries and dissolves into the grains to some extent so that no amorphous phase was found. For a further increase in the SiO2 addition, the ZrO2 matrix saturates with silicon, and hence, the excess amounts of SiO2 precipitate as glass pockets at multiple grain junctions. The critical SiO2 content, 0.3 wt%, in the present material coincides with earlier studies on changes in creep behavior [13] and in grain boundary electron conductivity [24], where the critical SiO2 content was 0.3 wt% for a grain size of 0.27 lm. 4.2. Comparison with earlier creep data The superplastic flow behavior of fine-grained Y-TZP has primarily been characterized by the n value defined by the following creep equation: e_ ¼ Arn d p expðQ=RT Þ;
ð3Þ
where e_ is the steady-state strain rate, d is the grain size, p is the grain size exponent, Q is the apparent activation energy, R is the gas constant, T is the absolute temperature and A is a constant. Earlier studies have noted that the stress–strain rate relationship of Y-TZP strongly depends on a small amount of impurities such as SiO2 and Al2 O3 [9–15,25–28]. For high-purity materials with
a total impurity content less than 0.1 wt%, the n value varies from 2.0 to P 3.0 with decreasing stress, whereas for low-purity materials, it takes 2.0 for all the examined stresses. In contrast, the present study shows that the SiO2 -doped 3Y-TZP exhibits a sigmoidal stress– strain rate relationship even for 0.3–2.5 wt% SiO2 . The present data for undoped and 0.07 wt%-SiO2 doped 3Y-TZP are compared with earlier ones for finegrained 2Y- or 3Y-TZP in Fig. 7, where the creep behavior of high- and low-purity materials has been examined in grain sizes of 0.3–0.4 lm [9,11–15,29]. An inspection of Fig. 7 clearly reveals that the trend of the increasing n value with decreasing stress is inherent for fine-grained Y-TZP. For example, examination by Hiraga et al. [13] has noted that, for 0–0.3 wt%-SiO2 -doped 3Y-TZP, the n value steeply decreases from 3.9 to 1.9 with SiO2 addition, but for 0.3–2.5 wt%-SiO2 -doped ones, an n value of 1.9 is independent of the SiO2 addition. Their data obtained at 10–40 MPa, however, locate in the transition stresses from Regions H to I for the present loading. Likewise, other data show a trend similar to the present data. At low stresses, the creep rates are higher in low-purity Y-TZP than in high-purity one as shown in Fig. 7 [9,11– 15,29]. However, a small n value close to unity did not appear in the earlier studies. This discrepancy may be related to the following reasons. First, Region L with n 1:3 shifts toward a lower stress region with SiO2 addition, and hence, it does not appear in the limited stress region of P 3 MPa in earlier studies. Second, as noted in the recent studies on undoped 3Y-TZP [17,18], a 10-1
Purity High Low
10
•
T (K) Present Data Hiraga et al. [13] Hines et al. [15] Wakai [29] Lakki et al. [14] Carry [10]
-2
Creep Rate, ε / s-1
3360
10-3
1673 3Y 1673 3Y 1673 2Y 1623 3Y 1623 2Y 1623 2-3Y
2.0
10-4 10-5 5.0
10-6 10-7 10-8
1.3
1
10 100 True Stress, σ / MPa
300
Fig. 7. Creep data for high- and low-purity Y-TZP plotted as a function of true stress, r. The data reported by Carry [10], Lakki et al. [14] and Wakai [29] were normalized at 1673 K using Eq. (3) with Q ¼ 580 kJ/mol [18].
K. Morita et al. / Acta Materialia 52 (2004) 3355–3364
4.3. The role of SiO2 addition in the creep behavior of fine-grained 3Y-TZP
Creep Rate, ε / s-1
10-2
•
Extrapolated
(a) Region H (50 MPa)
10-3
Region H (30 MPa)
10-4
Region I (10 MPa)
10-5
Region I (5 MPa)
10-6
Region L (1.5 MPa)
10-7
6 (b)
5
Stress Exponent, n
decrease in the creep rate due to grain growth during deformation leads to a spuriously large n value of 3.0 at low stresses. In earlier studies, however, the n values have been determined by ignoring the effect of grain growth and thereby a small n value close to unity may be masked. The present examinations indicate that the creep behavior of fine-grained 3Y-TZP is highly sensitive to the SiO2 -dependent microstructure, which correlated with the combination of grain size and the SiO2 content [20]. Even if the SiO2 content is constant, a decrease in the grain boundary area due to grain growth enhances the degree of silicon segregation or glass pocket precipitation [20,21]. These factors were slightly different from each other in earlier studies, resulting in the scatter of creep data and the variation in the n value as shown in Fig. 7. Nevertheless, the analysis of creep data showing a similar trend strongly supports the fact that the sigmoidal stress–strain rate relationship is inherent for finegrained Y-TZP, and the doped SiO2 enhances the creep rates at low stresses but decreases the creep rates at high stresses.
3361
Glass Pocket
4 3 Region I
2
Region H Region L
1
The SiO2 -dependent e_ rd and n values are summarized in Fig. 8(a) and (b), respectively. Although the n value in Region L is constant independently of the SiO2 content, the e_ rd and n values in the other regions sensitively change with the SiO2 content. The change in the e_ rd and n values simultaneously occurs with the SiO2 addition and completes with the clear occurrence of glass pocket, except for the slight increase in the creep rate with the SiO2 content. This SiO2 -dependent creep behavior indicates that the glass pockets precipitated at P 0.3 wt% SiO2 can slightly increase the creep rate by accelerating GBS process by the viscous flow of the glass phase [6], but does not change the rate-controlling mechanism of creep deformation. For the SiO2 -doped 3Y-TZP, only the silicon that segregates along boundaries and dissolves into ZrO2 grains seems to affect the rate-controlling mechanism. The influence of the silicon on the rate-controlling mechanism, however, seems to be different among Regions L, I and H. In Region L, the silicon does not change the rate-controlling mechanism but merely enhances the accommodation process. In Regions I and H, on the other hand, the silicon changes the rate-controlling mechanism. 4.4. Implications for rate-controlling mechanisms of SiO2 doped 3Y-TZP 4.4.1. Region L The small n value close to unity for the undoped 3YTZP can be explained from Nabarro–Herring (N–H)
0
0
0.5
1.0 1.5 2.0 SiO2 Content, wt%
2.5
3.0
Fig. 8. (a) The e_ rd and (b) n values in Regions L, I and H plotted as a function of the SiO2 content. The e_ rd values of 0.9–2.5 wt%-SiO2 doped materials at 50 MPa were evaluated from an extrapolation with n ¼ 1:8 in Fig. 5.
diffusion creep, in which the lattice diffusion of cations is the rate-controlling process [17,18]. An n value of 1.3 in Region L suggests that the deformation of SiO2 doped Y-TZP can also be ascribed to N–H creep. The increase in the creep rate for <0.3 wt% SiO2 implies an enhancement of lattice diffusion. It has been shown that SiO2 addition enhances not only sintering but also the grain growth of Y-TZP [30,31]. Since the rate of grain growth of Y-TZP is controlled by the lattice diffusion of cations [32], the silicon addition to the ZrO2 grains may accelerate the lattice diffusion of cations. For <0.3 wt% SiO2 , where the dissolution of silicon increases with the SiO2 addition, the enhanced lattice diffusivity of cations may result in the steep increase in the creep rate of Y-TZP. The slight increase in the creep rate for P 0.3 wt% SiO2 , however, cannot be ascribed to the enhanced lattice diffusivity by the silicon dissolution. This is because the glass pocket precipitation at P 0.3 wt% SiO2 suggests that the dissolution of silicon into the ZrO2 matrix is saturated. For ceramic materials with glass phases, the enhanced creep rate has often been related to
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the solution-precipitation model, where enhanced matter transport is assumed to occur through the glass phases [29]. If the matter transport in the glass pockets is the rate-controlling mechanism, the sigmoidal relationship between stress and strain rate should disappear with the occurrence of glass pocket at P 0.3 wt% SiO2 . This is contradictory to the results shown in Figs. 5 and 7, indicating that the solution-precipitation mechanism is not the rate-controlling mechanism of the SiO2 -doped 3Y-TZP. Diffusion creep in polycrystalline materials must be accompanied by GBS to preserve material continuity [33]. The GBS process is constrained by the surrounding grains at multiple grain junctions. According to the study on 5 wt%-SiO2 -doped ZrO2 [6], where about 10 vol% glass phase exists, the stress concentrations in the junctions exerted by GBS seems to be relived through the glass phase, resulting in accelerated plastic flow. In the present material, the nanometer-sized glass pockets precipitated at the junctions may also assist the relaxation of the stress concentrations by viscous flow of the glass phase. The slight increase in the creep rates for P 0.3 wt% SiO2 , therefore, may be ascribed to the accelerated GBS process caused by the increase in the volume fraction of the glass pockets. 4.4.2. Regions I and H The stress region with n 2:0 has generally been recognized that the deformation primarily occurs through the GBS mechanism, so that the creep rate is controlled by the rate of the accommodation process of stress concentrations exerted around multiple grain junctions. To explain the superplastic flow behavior of low-purity Y-TZP, several diffusion-related mechanisms have been proposed: (i) a solution-precipitation mechanism [29], (ii) the existence of an amorphous phase at elevated temperatures [28] and (iii) an interface-reaction controlled mechanism [10,15,25–27,34]. However, the present SiO2 -dependent creep behavior shown in Fig. 4 seems to be incompatible with these existing models. In earlier studies, it has been believed that grain boundary amorphous phase is essential for enhancing the superplastic flow [6,12,28,29]. For the solution-precipitation mechanism proposed by Wakai [29], deformation occurs by enhanced matter flow through the intergranular amorphous phase. Jimenez-Melendo et al. [28] have asserted that, at high temperatures, a grain boundary amorphous phase is formed along grain boundaries and hence creep rates were accelerated in low purity materials. However, the present microstructural examination demonstrated that no amorphous phase was formed along boundaries even at high temperatures. In contrast to their models, the creep rate is lowered by the SiO2 addition in Region H. In addition, their models predict an n value of 2.0 for the entire stress
region and cannot explain the sigmoidal e_ rd –r relationship of SiO2 -doped Y-TZP. To explain the flow mechanism of fine-grained Y-TZP without grain boundary amorphous phases, Chokshi and his colleagues [15,26,34] have proposed the operation of two sequential mechanisms: GBS at high stresses with n 2:0 and the interface-reaction controlled GBS at low stresses with n P 3:0. Superplastic deformation consists of the sequential processes of GBS, which involves nucleation and annihilation of defects at the interfaces, and its accommodation. At high stresses with n 2:0, GBS rapidly occurs so that the rate of deformation is controlled by the accommodation process of the predominant GBS mechanism, independent of the impurity content. At low stresses, GBS can be assumed to become slow enough to become the ratecontrolling process. Thus, for low-purity Y-TZP, the interface-reaction process would be enhanced by the presence of trace impurities, and hence, the stress region with n 2:0 extends to low stresses. For high-purity Y-TZP, on the other hand, the deformation occurs through the interface-reaction controlled GBS, providing n P 3:0 at low stresses. The interface-reaction controlled GBS mechanism appears to be compatible with the result of the mechanical spectroscopy by Lakki et al. [14]. Mechanical loss due to internal friction increases exponentially with temperature and the loss level is sufficiently higher in low-purity Y-TZP than that of high-purity one [14]. They have pointed out that since the internal friction can be attributed to GBS, GBS would more easily occur in low-purity Y-TZP. However, the interface-reaction controlled GBS mechanism cannot explain the creep rate decreasing with the SiO2 addition in Region H. In general, the high-temperature deformation is controlled by the rate of the accommodation process of the predominant GBS mechanism. For the SiO2 -doped 3Y-TZP, although the silicon preferentially segregated along the grain boundaries may indeed enhance GBS as noted by Lakki et al. [14], the SiO2 addition would decrease the rate of the accommodation process at high stresses. These results imply that, for the SiO2 -doped 3Y-TZP, other models seem to be necessary to explain the ratecontrolling mechanism in the superplastic region with n 2:0. The lattice diffusivity of cations seems to be enhanced by the dissolution of silicon into ZrO2 . In Region I, the creep rate of the low-purity materials is also enhanced due probably to the intervention of the enhanced lattice diffusivity, whereas in Region H, it is almost the same as or slightly slower than that of highpurity ones as shown in Figs. 4 and 6. At high stresses, diffusion-related mechanisms are unlikely for the accommodation process of stress concentrations exerted by GBS. The existence of intragranular dislocations also supports the fact that the stress concentrations cannot
K. Morita et al. / Acta Materialia 52 (2004) 3355–3364
be fully relaxed only by diffusional processes. At high stresses, dislocation-related mechanisms may also contribute to the accommodation process of GBS in the SiO2 -doped 3Y-TZP as in the case of high-purity Y-TZP [18]. For superplastic ceramics, dislocation-related mechanisms have been ruled out as the accommodation process of GBS because of the smaller flow stresses. Indeed, the applied stresses are about one order of magnitude smaller as compared with the yield stress of easy slip system in cubic ZrO2 single crystals [35]. In the recent studies on high-purity Y-TZP [18,36], however, we have demonstrated that the stress concentrations around multiple grain junctions exerted by GBS reach a factor of 14–25 [36]. The noticeably activated dislocation motion seen in Fig. 6(c) strongly suggests that the significant stress concentrations must reach a critical stress for dislocation nucleation even at the present loading and result in an intragranular dislocation motion. According to Parthasarathy and Hay [35], the relationship between dislocation density q and flow stress for the cubic ZrO2 single crystals can be expressed by the following equation: q ¼ 1:18 107 r2:36 :
ð4Þ
Assuming that a stress concentration factor of 10–20 occurs around multiple grain junctions, the q value can be estimated to be 1.2 1013 –6.1 1013 m2 for r ¼ 30 MPa, roughly corresponding to 1.2–5.9 dislocations per one grain. This predicted q value is also consistent with that observed at r ¼ 30 MPa in Fig. 6(c). Since GBS in a polycrystalline matrix cannot occur without a concomitant accommodation process for the stress concentrations, the rate of deformation is controlled by the recovery of the dislocations. For the SiO2 doped 3Y-TZP, the solute silicon may hinder the recovery process of dislocations within the grains or along the grain boundaries. The contribution of the dislocations to the accommodation process increases with an increase in stress [18], so that the influence of silicon may become noticeable with stress. For further discussion of the ratecontrolling mechanism in the superplastic region, the detailed determination of the p and Q values and a microstructural examination seem to be necessary.
5. Summary The correlation between the microstructure and tensile creep behavior of 3Y-TZP was examined as a function of SiO2 content. The results are summarized as follows: (1) For a small amount of SiO2 addition less than 0.3 wt%, which is the critical value for glass-pocket precipitation for a grain size of 0.3 lm, silicon seg-
3363
regates along the grain boundaries and dissolves into the grains to some extent. For the SiO2 addition of more than 0.3 wt%, the excess amount of SiO2 precipitates at multiple grain junctions. Although the grain boundaries of SiO2 -doped materials are enriched with silicon, no intragranular amorphous phase is found before and after deformation. (2) About 0–2.5 wt%-SiO2 -doped 3Y-TZP exhibits a sigmoidal stress-creep rate relationship, which correlates closely to a change in the microstructure due to the SiO2 addition. For <0.3 wt% SiO2 , the sigmoidal relationship changes drastically with the SiO2 addition. For P 0.3 wt% SiO2 , it is insensitive to an increase in the SiO2 content though the relationship gradually shifts toward the high strain rate region with an increase in the volume fraction of the glass pockets. (3) For SiO2 -doped materials, the glass pockets enhance the creep rate by accelerating the GBS process by viscous flow of the glass phase, but do not affect the rate-controlling mechanism. Only the silicon segregated along the grain boundaries and dissolved into grains affects the rate-controlling process. (4) In Region L with n 1:3, deformation can be attributed to Nabarro–Herring creep. The silicon dissolved into the ZrO2 grains may accelerate the lattice diffusivity of cations, resulting in the enhanced creep rates. In Region H with n 2:0, the solute silicon may change the rate-controlling mechanism by inhibiting the recovery of the intragranular dislocations, which contributes to the accommodation process of the predominant GBS mechanism.
References [1] Shackelford JF, Nicholson PS, Smeltzer WW. Am Ceram Soc Bull 1974;53:865. [2] Mecartney ML. J Am Ceram Soc 1987;70:54. [3] Hwang CMJ, Chen IW. J Am Ceram Soc 1990;73:1626. [4] Yoshizawa Y, Sakuma T. J Am Ceram Soc 1990;73:3069. [5] Kajihara K, Yoshizawa Y, Sakuma T. Scr Mater 1993;28:559. [6] Kajihara K, Yoshizawa Y, Sakuma T. Acta Mater 1995;43:1235. [7] Sharif AA, Mecartney ML. Acta Mater 2003;51:1633. [8] Ikuhara Y, Thavorniti P, Sakuma T. Acta Mater 1997;45:5275. [9] Hermansson T, Swan H, Dunlop G. In: With GD et al., editors. Euro-Ceramics. Engineering Ceramics, vol. 3. Elsevier Applied Science; 1989. p. 3.329. [10] Carry C. Mater Res Soc Symp Proc 1990;196:314. [11] Nauer M, Carry C. Scr Mater 1990;24:1459. [12] Gust M, Goo G, Wolfenstine J, Mecartney ML. J Am Ceram Soc 1993;76:1681. [13] Hiraga K, Yasuda HY, Sakka Y. Mater Sci Eng A 1997;234– 236:1026. [14] Lakki A, Schaller R, Nauer M, Carry C. Acta Mater 1993;41:2845. [15] Hines JA, Ikuhara Y, Chokshi AH, Sakuma T. Acta Mater 1998;46:5557. [16] Mendelson MI. J Am Ceram Soc 1969;52:443. [17] Morita K, Hiraga K. Scr Mater 2000;42:183.
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[18] Morita K, Hiraga K. Acta Mater 2002;50:1075. [19] Clarke DR. J Phys (Paris) 1985;46:51. [20] Morita K, Sakka Y, Hiraga K. Mater Mater Sci Forum 1999;304– 306:663. [21] Ikuhara Y, Nagai Y, Yamamoto T, Sakuma T. Mater Sci Forum 1999;304–306:525. [22] Aoki M, Chiang YM, Kosacki I, Lee JR, Tuller H, Liu Y. J Am Ceram Soc 1996;79:1169. [23] Gremillard L, Epicier T, Chevalier J, Fantozzi G. Acta Mater 2000;48:4647. [24] Uchikoshi T, Sakka Y, Hiraga K. J Electroceram 1999;47:3433. [25] Langdon TG. In: Bradt RC et al., editors. Plastic Deformation of Ceramics. New York: Plenum Press; 1995. p. 251. [26] Chokshi AH. Mater Sci Eng A 1993;166:119.
[27] Sato E, Morioka H, Kuribayashi K, Sundararaman D. J Mater Sci 1999;34:4511. [28] Jimenez-Melendo M, Domınguez-Rodorıguez A, Bravo-Le on A. J Am Ceram Soc 1998;81:2761. [29] Wakai F. Acta Mater 1994;42:1163. [30] Badwel SPS, Rajendran S. Solid State Ionics 1994;70/71:83. [31] G€ odickemeier M, Michel B, Orliukas A, Bohac P, Sasaki K, Gauckler L, et al. J Mater Res 1994;9:1228. [32] Zhao J, Ikuhara Y, Sakuma T. J Am Ceram Soc 1998;81: 2087. [33] Stevens RN. Philos Mag 1971;23:265. [34] Owen DM, Chokshi AH. Acta Mater 1998;46:667. [35] Parthasarathy TA, Hay RS. Acta Mater 1996;44:4663. [36] Morita K, Hiraga K. Philos Mag Lett 2001;81:311.
Effect of minor SiO2 addition on the creep behavior of superplastic tetragonal ZrO2 K. Morita *, K. Hiraga, B.-N. Kim National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan Received 2 December 2003; received in revised form 22 March 2004; accepted 24 March 2004 Available online 21 April 2004
Abstract A sigmoidal stress–creep rate relationship appearing in tetragonal ZrO2 drastically changes with SiO2 addition, where silicon segregates along the grain boundaries and dissolves into the ZrO2 grains. This change is completed with the precipitation of glass pockets though the creep rate gradually increases with increasing the volume fraction. The glass pockets merely increase the creep rate by accelerating grain boundary sliding (GBS) by the viscous flow of the glass phase. Although the silicon segregated along the grain boundaries may enhance GBS, the solute silicon seems to control the rate of creep. The solute silicon increases the rate of Nabarro–Herring creep by accelerating the lattice diffusivity of cations at low stresses but decreases the rate of GBS by inhibiting the accommodation process at high stresses. Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Ceramics; High-temperature deformation; Superplasticity; Tetragonal ZrO2
1. Introduction Silicate glasses have deliberately been added to ZrO2 ceramics because they can enhance sintering [1,2]. The addition of glasses is known to play an important role in the mechanical properties of ZrO2 ceramics, particularly at elevated temperatures [3–7]. The addition of pure SiO2 can enhance the superplasticity of fine-grained Y2 O3 -stabilized tetragonal ZrO2 polycrystal (Y-TZP) [6,7]; e.g., 5 wt%-SiO2 -doped Y-TZP exhibits a large tensile ductility exceeding 1000% at 1673 K [6]. The enhanced superplasticity of SiO2 doped Y-TZP has been attributed to plastic flow accelerated through glass pockets [6] and a change in the chemical bonding state of grain boundaries caused by the segregation of silicon [8]. The former factor can enhance the accommodation process of stress concentrations exerted at multiple grain junctions by grain boundary sliding (GBS) and the latter factor can strengthen the bonding of the grain boundaries. *
Corresponding author. Tel.: +81-29-859-2537; fax: +81-29-8592501. E-mail address: [email protected] (K. Morita).
Earlier studies [9–15] have noted that the superplastic flow behavior of Y-TZP is sensitive to the existence of residual impurities. This is because the impurities significantly affect GBS that has generally been accepted as the primary flow mechanism. For a large amount of SiO2 addition, both factors simultaneously work on the superplastic flow and hence it is hard to examine the respective roles of SiO2 addition. For understanding the exact role of SiO2 addition in the superplasticity of Y-TZP, the deformation behavior should be examined over a wide range of SiO2 addition with and without glass phase precipitation. The present study was therefore performed to examine the effect of SiO2 on the creep behavior of finegrained Y-TZP as a function of the SiO2 content. The creep behavior was correlated with the microstructure varied with the SiO2 content ranging from 0 to 2.5 wt%.
2. Experimental procedures The materials used were prepared from a high-purity 3Y-TZP powder containing 3 mol% Y2 O3 (Tosoh Co., Ltd.,: 50 SiO2 , <50 Al2 O3 , <50 Fe2 O3 , 220 NaO in wt
1359-6454/$30.00 Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2004.03.033
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ppm) and pure SiO2 (Nissan Chemical Industry Co., Ltd.) using a method employed by Kajihara et al. [5]. The 3Y-TZP powders doped with 0–2.5 wt% SiO2 were cold isostatically pressed at about 200 MPa, sintered at 1673 K for 3 h in air and cooled in the furnace. The sintered materials had a relative density higher than 98%. The initial grain size, d0 , defined as 1.56 times the average intercept lengths of the grains [16] was about 0.35 lm. From the sintered bodies, dog-bone-shaped flat tensile specimens were machined with a gauge portion of t w l 3- 3- 20 mm. Tensile creep tests were conducted at 1673 K in air and at initial stresses of 0.8–80 MPa. We corrected the interference arising from the deformation around the grip sections and grain growth using a method described elsewhere [17,18]. Briefly, true displacement was directly measured by monitoring the length between the targets made at both ends of the gauge portion. The monitored creep rates, e_ m , can be corrected for instantaneous stress as e_ r ¼ e_ m expðnea Þ;
ð1Þ
where ea is the arbitrary strain. When grain growth occurs during deformation, Eq. (1) can be corrected for the instantaneous grain size, da , at ea as e_ rd ¼ e_ m expðnea Þðd0 =da Þp :
ð2Þ
The as-sintered and deformed microstructures were examined by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Before the SEM observations, the specimens were mechanically polished and thermally etched at 1473 K for 30 min. For the TEM observation of the as-sintered microstructures, several specimens were quenched from 1673 K by di-
rectly dropping from the furnace into water so that they must preserve the microstructure at high-temperature. Sheets with a thickness of 500 lm were cut with a lowspeed diamond cutter, mechanically polished to less than 100 lm in thickness and further thinned with an Ar ion-milling machine. The energy dispersive X-ray spectroscopy (EDS) was performed at areas thin enough to obtain lattice image using a JEOL-2010F with a probe size of about 1.0 nm.
3. Experimental results 3.1. As-sintered microstructure As shown in Fig. 1, the microstructure of the as-sintered materials depends on the SiO2 content and can be classified into two groups. For a small amount of the SiO2 addition less than 0.3 wt%, no glass phase is found at any grain junction, but beyond this, nanometer-sized glass pockets are apparently observed to precipitate at rounded multiple grain junctions as indicated by the triangles. At P 0.3 wt% SiO2 , the volume fraction of the glass pockets monotonously increases with the SiO2 content. Further details of the microstructure were examined using high-resolution TEM and EDS in Fig. 2. EDS analysis clearly revealed that the nanometer-sized glass pockets in the 0.3–2.5 wt%-SiO2 -doped materials consisted mainly of SiO2 . For the undoped materials, the segregation of yttrium was found along the grain boundaries, whereas for the SiO2 -doped materials, silicon was found to segregate along the grain boundaries
Fig. 1. Bright-field TEM images of the as-sintered 3Y-TZP doped with (a) 0, (b) 0.07, (c) 0.3 and (d) 2.5 wt% SiO2 .
K. Morita et al. / Acta Materialia 52 (2004) 3355–3364
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On the other hand, EDS spectra obtained from the grain interior revealed a trace of silicon in addition to zirconium and yttrium (Fig. 2(c)). Since the EDS measurements were carried out at sufficient distances, at least >20 nm, from the grain boundaries and glass pockets, the interference from those areas can be negligible even if beam scattering occurred through thin foils. This shows that a small amount of silicon also dissolves into ZrO2 grains. The level of silicon along the grain boundaries and within the grains was deduced from the X-ray intensity of the Si–K peak in the EDS spectra. As shown in Fig. 3, the level of silicon depends on the SiO2 content. For <0.3 wt% SiO2 , the level steeply increases with the SiO2 addition, whereas for P 0.3 wt% SiO2 , it is almost constant independently of the SiO2 content. The saturation value is compatible with earlier data for 5 wt%-SiO2 -doped 2.5Y-TZP [8]. Although the level of silicon along the grain boundaries is apparently higher than that of the grain interiors, it tends to be lower in the water-quenched materials than in the furnace-cooled ones. 3.2. Tensile creep behavior The creep rate curve of the present materials strongly depends on the correction of grain size as shown in Fig. 4. For the e_ m and e_ r curves, the creep rates gradually change with e and no well-defined steady-state region appears. For the e_ rd curves, on the other hand, a well-defined steady-state region appears at e P 0:03. It is noticeable that the discrepancy between the e_ r and e_ rd curves increases with e and it becomes more prominent at lower stresses; it exceeds a factor of 2 at e P 0:1 for 3 MPa. Thus, the occurrence of grain growth strongly affects the creep behavior of fine-grained 3Y-TZP. Since 0.20
Fig. 2. (a) High-resolution TEM image of a multiple grain junction filled with the glass phase and typical EDS spectra obtained from (b) grain boundaries and (c) grain interiors for the water-quenched 3YTZP doped with 0.3 wt% SiO2 .
in addition to yttrium as shown in Fig. 2(b). In the present materials, however, the lattice fringes of adjacent grains directly intersect at the boundaries and no amorphous phase was found along the grain boundaries as shown in Fig. 2(a). If an amorphous phase exists along grain boundaries, the dihedral angle at glass pockets must be less than 10° [19,20]. The large dihedral angles of 60° strongly support the fact that the grain boundaries of the SiO2 -doped materials are not wetted by amorphous phases even at high-temperatures [8].
Intensity Ratio, Si/Zr
furnace- watercooled quenched
0.15
Grain Boundary Grain Interior
Ikuhara et al. [8]
0.10
0.05
0
0
0.5
1.0
1.5
2.0
2.5
3.0
5.0
SiO2 Content, wt% Fig. 3. The X-ray intensity of the Si–K peak normalized by the Zr–L peak, Si/Zr, plotted as a function of the SiO2 content. The circles and squares represent the data of the furnace-cooled and water-quenched specimens, respectively. For comparison, earlier data for 5 wt%-SiO2 doped 2.5Y-TZP [8] are also plotted.
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(a) 3Y-TZP + 0.3wt%SiO2
the creep behavior should be characterized in the steadystate region, the grain size correction is necessary for the SiO2 -doped 3Y-TZP, as noted in the recent study for high-purity 3Y-TZP [17].
1673 K - 3 MPa
Creep Rate, ε / s-1
10-5
•
•
εσd •
ε•m εσ
10-6
(b) 3Y-TZP + 0.3wt%SiO2
1673 K - 30 MPa
-3
•
10
Creep Rate, ε / s-1
3.3. Stress–strain rate relationship
εm •
εσd
•
•
εσ
10-4 0
0.05
0.1
True Strain, ε
0.15
0.2
Fig. 4. Creep rate curves plotted as a function of true strain, e, for 0.3 wt%-SiO2 -doped 3Y-TZP obtained at 1673 K and at (a) 3 and (b) 30 MPa.
The steady-state creep rates, e_ rd , for 3Y-TZP without and with glass pockets are plotted as a function of true stress, r, in Fig. 5(a) and (b), respectively. The e_ rd –r relationship shows three prominent features. First, the e_ rd –r relationship is divided into three categories; a high-stress region (Region H) with n 1:8–2:7, an intermediate-stress region (Region I) with n 2:5–5:0 and a low-stress region (Region L) with n 1:3. Thus, the e_ rd –r relationship as a whole exhibits a sigmoidal feature in all the SiO2 contents examined though the sigmoidal feature becomes less noticeable for P 0.3 wt% SiO2 . Second, the sigmoidal e_ rd –r relationship is not a simple function of the SiO2 content. For <0.3 wt% SiO2 , the sigmoidal e_ rd –r relationship sensitively changes with the SiO2 addition (Fig. 5(a)), whereas for P 0.3 wt% SiO2 , the general feature is insensitive to the SiO2 addition though the creep rates gradually increase with the SiO2 content (Fig. 5(b)). Third, the effect of the SiO2 addition is different among Regions H, I and L. The n value is dependent on
10-1
Corrected Creep Rate, εσd / s-1
(a) without glass pocket
•
SiO2 (wt%) 0 [17]
10-2 10
(b) With glass pocket Kajihara et al. [6] 2.5Y-TZP+5wt%SiO2 D0 = 0.25µm
2.7
0.02 0.07
-3
1.8
10-4
1.8
2.5
10-5 1.3
2.5
5.0
0.07 0.3 0.9 2.5
10-6 1.3
10-7 Region L
10
Region I
Region H
Region L
SiO2 (wt%)
Region I
Region H
-8
1
10 100 True Stress, σ / MPa
1
10 100 True Stress, σ / MPa
Fig. 5. Corrected creep rates, e_ rd , plotted as a function of true stress, r, for 3Y-TZP doped with (a) 0–0.07 and (b) 0.07–2.5 wt% SiO2 . For comparison, earlier data for 5 wt%-SiO2 -doped 2.5Y-TZP [6] are also plotted. The open and closed symbols represent the creep data for the materials without and with glass pockets, respectively.
K. Morita et al. / Acta Materialia 52 (2004) 3355–3364
3359
the SiO2 content in Regions H and I, but independent of the content in Region L. In Regions H, for example, the SiO2 addition tends to decrease the creep rate and steeply decreases the n value from 2.7 to 1.8 up to 0.3 wt%, but beyond this, it slightly enhances the creep rate with maintaining n 1:8. In Region L, on the other hand, the SiO2 addition enhances the creep rate but does not change the n value of 1.3. According to Gust et al. [12], the addition of a small amount of barium silicate glass appears to increase the shear modulus, G, of 3Y-TZP. If this is also the case for the SiO2 -doped 3Y-TZP, the SiO2 -dependent creep behavior shown in Fig. 5 may be related to a change in the G value due to glass addition. However, the change in the G value is 4% [12] and negligible. In addition, the e_ rd –r relationship is not a simple function of the SiO2 content. Thus, the general feature of the SiO2 -dependent e_ rd –r relationship cannot be ascribed to a change in the G value accompanying the SiO2 addition. 3.4. Deformed microstructure Fig. 6 shows an example of a deformed microstructure. After deformation, the specimen was quickly cooled under loading to preserve the microstructures developed during deformation. Although the deformed materials almost retained the equiaxed grain shapes, strain contrasts appeared to be more pronounced after deformation. The co-segregation of silicon and yttrium was also detected along grain boundaries, but no intergranular amorphous phase was found even after deformation, as shown in Fig. 6(b). 1 Dark-field TEM observation reveals densely aligned intragranular dislocations, which seem to nucleate around multiple grain junctions indicated by an arrow in Fig. 6(c). Since these dislocation substructures were not observed before deformation, the dislocations are developed during deformation. 4. Discussion 4.1. Microstructural change due to SiO2 addition In earlier studies on glass-doped Y-TZP [5,8–14], it has been believed that the solubility of silicon into ZrO2 grains was highly limited, and hence, it formed an intergranular amorphous film with a thickness of about 1.0 nm. In contrast, the present study shows that a small amount of silicon dissolves into the ZrO2 grains. In 1
Glass phases precipitated at multiple grain junctions were observed to penetrate into two grain junctions only around the ahead of cracklike cavities, which are formed at e > 0:8 [20]. However, such crack-like cavities were not observed at e < 0:15, where the e_ rd values were determined, and no grain boundary amorphous phase was found in the present study.
Fig. 6. (a) Bright-field, (b) high-resolution and (c) dark-field TEM images for 2.5 wt%-SiO2 -doped 3Y-TZP after deformation up to e 0:18 at 1673 K and at 30 MPa.
addition, although the grain boundaries of the SiO2 doped materials are enriched with silicon, no amorphous phase occurs along the grain boundaries even after deformation. The present results are consistent with the earlier studies by Ikuhara et al. [8,21] and Aoki et al. [22], showing the segregation of silicon but no
K. Morita et al. / Acta Materialia 52 (2004) 3355–3364
amorphous phase along grain boundaries and the dissolution of silicon into ZrO2 grains. The level of silicon segregation along the grain boundaries tends to be higher in the furnace-cooled materials than in the water-quenched ones. According to the TEM in-situ heating experiment by Ikuhara et al. [21], at high temperatures, a small amount of SiO2 is soluble into the ZrO2 grains in the vicinity of the grain boundaries, but during cooling, the dissolved SiO2 precipitates from the super-saturated SiO2 –ZrO2 solid solution. This result suggests that the higher silicon segregation in the furnace-cooled materials may be caused by the precipitation of SiO2 , which dissolved into the ZrO2 grains at high temperatures. The recent study of Gremillard et al. [23] has noted that no silicon enrichment was observed along grain boundaries and no detectable silicon was found in grain interiors in 2.5 wt%-SiO2 -doped 3Y-TZP. Although the reason for the discrepancy is not clear at the present time, the conclusion clearly contradicts to the present microstructural examination, showing that no glass pocket was found for SiO2 additions less than 0.3 wt%. If silicon is insoluble into the ZrO2 matrix, glass pockets should precipitate simultaneously with the SiO2 addition. In addition, we confirmed that a decrease in the grain boundary area due to grain growth enhanced the precipitation of the glass pockets [20]. As a result, the change in the microstructure accompanying the SiO2 addition can be explained as follows. For a small amount of the SiO2 addition less than 0.3 wt%, the doped SiO2 preferentially segregates along the grain boundaries and dissolves into the grains to some extent so that no amorphous phase was found. For a further increase in the SiO2 addition, the ZrO2 matrix saturates with silicon, and hence, the excess amounts of SiO2 precipitate as glass pockets at multiple grain junctions. The critical SiO2 content, 0.3 wt%, in the present material coincides with earlier studies on changes in creep behavior [13] and in grain boundary electron conductivity [24], where the critical SiO2 content was 0.3 wt% for a grain size of 0.27 lm. 4.2. Comparison with earlier creep data The superplastic flow behavior of fine-grained Y-TZP has primarily been characterized by the n value defined by the following creep equation: e_ ¼ Arn d p expðQ=RT Þ;
ð3Þ
where e_ is the steady-state strain rate, d is the grain size, p is the grain size exponent, Q is the apparent activation energy, R is the gas constant, T is the absolute temperature and A is a constant. Earlier studies have noted that the stress–strain rate relationship of Y-TZP strongly depends on a small amount of impurities such as SiO2 and Al2 O3 [9–15,25–28]. For high-purity materials with
a total impurity content less than 0.1 wt%, the n value varies from 2.0 to P 3.0 with decreasing stress, whereas for low-purity materials, it takes 2.0 for all the examined stresses. In contrast, the present study shows that the SiO2 -doped 3Y-TZP exhibits a sigmoidal stress– strain rate relationship even for 0.3–2.5 wt% SiO2 . The present data for undoped and 0.07 wt%-SiO2 doped 3Y-TZP are compared with earlier ones for finegrained 2Y- or 3Y-TZP in Fig. 7, where the creep behavior of high- and low-purity materials has been examined in grain sizes of 0.3–0.4 lm [9,11–15,29]. An inspection of Fig. 7 clearly reveals that the trend of the increasing n value with decreasing stress is inherent for fine-grained Y-TZP. For example, examination by Hiraga et al. [13] has noted that, for 0–0.3 wt%-SiO2 -doped 3Y-TZP, the n value steeply decreases from 3.9 to 1.9 with SiO2 addition, but for 0.3–2.5 wt%-SiO2 -doped ones, an n value of 1.9 is independent of the SiO2 addition. Their data obtained at 10–40 MPa, however, locate in the transition stresses from Regions H to I for the present loading. Likewise, other data show a trend similar to the present data. At low stresses, the creep rates are higher in low-purity Y-TZP than in high-purity one as shown in Fig. 7 [9,11– 15,29]. However, a small n value close to unity did not appear in the earlier studies. This discrepancy may be related to the following reasons. First, Region L with n 1:3 shifts toward a lower stress region with SiO2 addition, and hence, it does not appear in the limited stress region of P 3 MPa in earlier studies. Second, as noted in the recent studies on undoped 3Y-TZP [17,18], a 10-1
Purity High Low
10
•
T (K) Present Data Hiraga et al. [13] Hines et al. [15] Wakai [29] Lakki et al. [14] Carry [10]
-2
Creep Rate, ε / s-1
3360
10-3
1673 3Y 1673 3Y 1673 2Y 1623 3Y 1623 2Y 1623 2-3Y
2.0
10-4 10-5 5.0
10-6 10-7 10-8
1.3
1
10 100 True Stress, σ / MPa
300
Fig. 7. Creep data for high- and low-purity Y-TZP plotted as a function of true stress, r. The data reported by Carry [10], Lakki et al. [14] and Wakai [29] were normalized at 1673 K using Eq. (3) with Q ¼ 580 kJ/mol [18].
K. Morita et al. / Acta Materialia 52 (2004) 3355–3364
4.3. The role of SiO2 addition in the creep behavior of fine-grained 3Y-TZP
Creep Rate, ε / s-1
10-2
•
Extrapolated
(a) Region H (50 MPa)
10-3
Region H (30 MPa)
10-4
Region I (10 MPa)
10-5
Region I (5 MPa)
10-6
Region L (1.5 MPa)
10-7
6 (b)
5
Stress Exponent, n
decrease in the creep rate due to grain growth during deformation leads to a spuriously large n value of 3.0 at low stresses. In earlier studies, however, the n values have been determined by ignoring the effect of grain growth and thereby a small n value close to unity may be masked. The present examinations indicate that the creep behavior of fine-grained 3Y-TZP is highly sensitive to the SiO2 -dependent microstructure, which correlated with the combination of grain size and the SiO2 content [20]. Even if the SiO2 content is constant, a decrease in the grain boundary area due to grain growth enhances the degree of silicon segregation or glass pocket precipitation [20,21]. These factors were slightly different from each other in earlier studies, resulting in the scatter of creep data and the variation in the n value as shown in Fig. 7. Nevertheless, the analysis of creep data showing a similar trend strongly supports the fact that the sigmoidal stress–strain rate relationship is inherent for finegrained Y-TZP, and the doped SiO2 enhances the creep rates at low stresses but decreases the creep rates at high stresses.
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Glass Pocket
4 3 Region I
2
Region H Region L
1
The SiO2 -dependent e_ rd and n values are summarized in Fig. 8(a) and (b), respectively. Although the n value in Region L is constant independently of the SiO2 content, the e_ rd and n values in the other regions sensitively change with the SiO2 content. The change in the e_ rd and n values simultaneously occurs with the SiO2 addition and completes with the clear occurrence of glass pocket, except for the slight increase in the creep rate with the SiO2 content. This SiO2 -dependent creep behavior indicates that the glass pockets precipitated at P 0.3 wt% SiO2 can slightly increase the creep rate by accelerating GBS process by the viscous flow of the glass phase [6], but does not change the rate-controlling mechanism of creep deformation. For the SiO2 -doped 3Y-TZP, only the silicon that segregates along boundaries and dissolves into ZrO2 grains seems to affect the rate-controlling mechanism. The influence of the silicon on the rate-controlling mechanism, however, seems to be different among Regions L, I and H. In Region L, the silicon does not change the rate-controlling mechanism but merely enhances the accommodation process. In Regions I and H, on the other hand, the silicon changes the rate-controlling mechanism. 4.4. Implications for rate-controlling mechanisms of SiO2 doped 3Y-TZP 4.4.1. Region L The small n value close to unity for the undoped 3YTZP can be explained from Nabarro–Herring (N–H)
0
0
0.5
1.0 1.5 2.0 SiO2 Content, wt%
2.5
3.0
Fig. 8. (a) The e_ rd and (b) n values in Regions L, I and H plotted as a function of the SiO2 content. The e_ rd values of 0.9–2.5 wt%-SiO2 doped materials at 50 MPa were evaluated from an extrapolation with n ¼ 1:8 in Fig. 5.
diffusion creep, in which the lattice diffusion of cations is the rate-controlling process [17,18]. An n value of 1.3 in Region L suggests that the deformation of SiO2 doped Y-TZP can also be ascribed to N–H creep. The increase in the creep rate for <0.3 wt% SiO2 implies an enhancement of lattice diffusion. It has been shown that SiO2 addition enhances not only sintering but also the grain growth of Y-TZP [30,31]. Since the rate of grain growth of Y-TZP is controlled by the lattice diffusion of cations [32], the silicon addition to the ZrO2 grains may accelerate the lattice diffusion of cations. For <0.3 wt% SiO2 , where the dissolution of silicon increases with the SiO2 addition, the enhanced lattice diffusivity of cations may result in the steep increase in the creep rate of Y-TZP. The slight increase in the creep rate for P 0.3 wt% SiO2 , however, cannot be ascribed to the enhanced lattice diffusivity by the silicon dissolution. This is because the glass pocket precipitation at P 0.3 wt% SiO2 suggests that the dissolution of silicon into the ZrO2 matrix is saturated. For ceramic materials with glass phases, the enhanced creep rate has often been related to
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the solution-precipitation model, where enhanced matter transport is assumed to occur through the glass phases [29]. If the matter transport in the glass pockets is the rate-controlling mechanism, the sigmoidal relationship between stress and strain rate should disappear with the occurrence of glass pocket at P 0.3 wt% SiO2 . This is contradictory to the results shown in Figs. 5 and 7, indicating that the solution-precipitation mechanism is not the rate-controlling mechanism of the SiO2 -doped 3Y-TZP. Diffusion creep in polycrystalline materials must be accompanied by GBS to preserve material continuity [33]. The GBS process is constrained by the surrounding grains at multiple grain junctions. According to the study on 5 wt%-SiO2 -doped ZrO2 [6], where about 10 vol% glass phase exists, the stress concentrations in the junctions exerted by GBS seems to be relived through the glass phase, resulting in accelerated plastic flow. In the present material, the nanometer-sized glass pockets precipitated at the junctions may also assist the relaxation of the stress concentrations by viscous flow of the glass phase. The slight increase in the creep rates for P 0.3 wt% SiO2 , therefore, may be ascribed to the accelerated GBS process caused by the increase in the volume fraction of the glass pockets. 4.4.2. Regions I and H The stress region with n 2:0 has generally been recognized that the deformation primarily occurs through the GBS mechanism, so that the creep rate is controlled by the rate of the accommodation process of stress concentrations exerted around multiple grain junctions. To explain the superplastic flow behavior of low-purity Y-TZP, several diffusion-related mechanisms have been proposed: (i) a solution-precipitation mechanism [29], (ii) the existence of an amorphous phase at elevated temperatures [28] and (iii) an interface-reaction controlled mechanism [10,15,25–27,34]. However, the present SiO2 -dependent creep behavior shown in Fig. 4 seems to be incompatible with these existing models. In earlier studies, it has been believed that grain boundary amorphous phase is essential for enhancing the superplastic flow [6,12,28,29]. For the solution-precipitation mechanism proposed by Wakai [29], deformation occurs by enhanced matter flow through the intergranular amorphous phase. Jimenez-Melendo et al. [28] have asserted that, at high temperatures, a grain boundary amorphous phase is formed along grain boundaries and hence creep rates were accelerated in low purity materials. However, the present microstructural examination demonstrated that no amorphous phase was formed along boundaries even at high temperatures. In contrast to their models, the creep rate is lowered by the SiO2 addition in Region H. In addition, their models predict an n value of 2.0 for the entire stress
region and cannot explain the sigmoidal e_ rd –r relationship of SiO2 -doped Y-TZP. To explain the flow mechanism of fine-grained Y-TZP without grain boundary amorphous phases, Chokshi and his colleagues [15,26,34] have proposed the operation of two sequential mechanisms: GBS at high stresses with n 2:0 and the interface-reaction controlled GBS at low stresses with n P 3:0. Superplastic deformation consists of the sequential processes of GBS, which involves nucleation and annihilation of defects at the interfaces, and its accommodation. At high stresses with n 2:0, GBS rapidly occurs so that the rate of deformation is controlled by the accommodation process of the predominant GBS mechanism, independent of the impurity content. At low stresses, GBS can be assumed to become slow enough to become the ratecontrolling process. Thus, for low-purity Y-TZP, the interface-reaction process would be enhanced by the presence of trace impurities, and hence, the stress region with n 2:0 extends to low stresses. For high-purity Y-TZP, on the other hand, the deformation occurs through the interface-reaction controlled GBS, providing n P 3:0 at low stresses. The interface-reaction controlled GBS mechanism appears to be compatible with the result of the mechanical spectroscopy by Lakki et al. [14]. Mechanical loss due to internal friction increases exponentially with temperature and the loss level is sufficiently higher in low-purity Y-TZP than that of high-purity one [14]. They have pointed out that since the internal friction can be attributed to GBS, GBS would more easily occur in low-purity Y-TZP. However, the interface-reaction controlled GBS mechanism cannot explain the creep rate decreasing with the SiO2 addition in Region H. In general, the high-temperature deformation is controlled by the rate of the accommodation process of the predominant GBS mechanism. For the SiO2 -doped 3Y-TZP, although the silicon preferentially segregated along the grain boundaries may indeed enhance GBS as noted by Lakki et al. [14], the SiO2 addition would decrease the rate of the accommodation process at high stresses. These results imply that, for the SiO2 -doped 3Y-TZP, other models seem to be necessary to explain the ratecontrolling mechanism in the superplastic region with n 2:0. The lattice diffusivity of cations seems to be enhanced by the dissolution of silicon into ZrO2 . In Region I, the creep rate of the low-purity materials is also enhanced due probably to the intervention of the enhanced lattice diffusivity, whereas in Region H, it is almost the same as or slightly slower than that of highpurity ones as shown in Figs. 4 and 6. At high stresses, diffusion-related mechanisms are unlikely for the accommodation process of stress concentrations exerted by GBS. The existence of intragranular dislocations also supports the fact that the stress concentrations cannot
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be fully relaxed only by diffusional processes. At high stresses, dislocation-related mechanisms may also contribute to the accommodation process of GBS in the SiO2 -doped 3Y-TZP as in the case of high-purity Y-TZP [18]. For superplastic ceramics, dislocation-related mechanisms have been ruled out as the accommodation process of GBS because of the smaller flow stresses. Indeed, the applied stresses are about one order of magnitude smaller as compared with the yield stress of easy slip system in cubic ZrO2 single crystals [35]. In the recent studies on high-purity Y-TZP [18,36], however, we have demonstrated that the stress concentrations around multiple grain junctions exerted by GBS reach a factor of 14–25 [36]. The noticeably activated dislocation motion seen in Fig. 6(c) strongly suggests that the significant stress concentrations must reach a critical stress for dislocation nucleation even at the present loading and result in an intragranular dislocation motion. According to Parthasarathy and Hay [35], the relationship between dislocation density q and flow stress for the cubic ZrO2 single crystals can be expressed by the following equation: q ¼ 1:18 107 r2:36 :
ð4Þ
Assuming that a stress concentration factor of 10–20 occurs around multiple grain junctions, the q value can be estimated to be 1.2 1013 –6.1 1013 m2 for r ¼ 30 MPa, roughly corresponding to 1.2–5.9 dislocations per one grain. This predicted q value is also consistent with that observed at r ¼ 30 MPa in Fig. 6(c). Since GBS in a polycrystalline matrix cannot occur without a concomitant accommodation process for the stress concentrations, the rate of deformation is controlled by the recovery of the dislocations. For the SiO2 doped 3Y-TZP, the solute silicon may hinder the recovery process of dislocations within the grains or along the grain boundaries. The contribution of the dislocations to the accommodation process increases with an increase in stress [18], so that the influence of silicon may become noticeable with stress. For further discussion of the ratecontrolling mechanism in the superplastic region, the detailed determination of the p and Q values and a microstructural examination seem to be necessary.
5. Summary The correlation between the microstructure and tensile creep behavior of 3Y-TZP was examined as a function of SiO2 content. The results are summarized as follows: (1) For a small amount of SiO2 addition less than 0.3 wt%, which is the critical value for glass-pocket precipitation for a grain size of 0.3 lm, silicon seg-
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regates along the grain boundaries and dissolves into the grains to some extent. For the SiO2 addition of more than 0.3 wt%, the excess amount of SiO2 precipitates at multiple grain junctions. Although the grain boundaries of SiO2 -doped materials are enriched with silicon, no intragranular amorphous phase is found before and after deformation. (2) About 0–2.5 wt%-SiO2 -doped 3Y-TZP exhibits a sigmoidal stress-creep rate relationship, which correlates closely to a change in the microstructure due to the SiO2 addition. For <0.3 wt% SiO2 , the sigmoidal relationship changes drastically with the SiO2 addition. For P 0.3 wt% SiO2 , it is insensitive to an increase in the SiO2 content though the relationship gradually shifts toward the high strain rate region with an increase in the volume fraction of the glass pockets. (3) For SiO2 -doped materials, the glass pockets enhance the creep rate by accelerating the GBS process by viscous flow of the glass phase, but do not affect the rate-controlling mechanism. Only the silicon segregated along the grain boundaries and dissolved into grains affects the rate-controlling process. (4) In Region L with n 1:3, deformation can be attributed to Nabarro–Herring creep. The silicon dissolved into the ZrO2 grains may accelerate the lattice diffusivity of cations, resulting in the enhanced creep rates. In Region H with n 2:0, the solute silicon may change the rate-controlling mechanism by inhibiting the recovery of the intragranular dislocations, which contributes to the accommodation process of the predominant GBS mechanism.
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