High Strain in (K,Na)NbO3-Based Lead-Free Piezoelectric Fibers

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High Strain in (K,Na)NbO3‑Based Lead-Free Piezoelectric Fibers Francesca Bortolani,† Adolfo del Campo,‡ José F. Fernandez,‡ Frank Clemens,† and Fernando Rubio-Marcos*,‡ †

EMPA-Laboratory for High Performance Ceramics, Ueberlandstrasse 129, 8600 Dübendorf, Switzerland Electroceramic Department, Instituto de Cerámica y Vidrio, CSIC, Kelsen 5, 28049 Madrid, Spain

‡

S Supporting Information *

ABSTRACT: Until now, lead zirconate titanate (PZT)-based ceramics are the most widely used in piezoelectric devices. However, the use of lead is being avoided due to its toxicity and environmental risks. Indeed, the attention has been moved to lead-free ceramics, especially on potassium sodium niobate (KNN)-based materials, due to growing environmental concerns. These materials are technologically interesting. For applications such as actuators, an electromechanical coupling providing high strain with high force, e.g, fuel injection, ultrasonic motor, etc., is required, Moreover, in the current context, the new technologies evolve toward the miniaturization of the conventional electronic devices. Herein, we have developed microfiber ceramics of KNN-based composition, which yield a high strain value with Smax as high as 0.17% at 3 kV mm−1. According to our results, this phenomenon can be explained by an extrinsic effect that favors the internal relaxation of the system. To reach this breakthrough, a sintering mechanism has been established, which allows for correlating the extrinsic factors of the system with electromechanical properties of the ceramic fibers. We believe that the general strategy and design principles described in this study will open new avenues in developing of (K,Na)NbO3-based lead-free piezoelectric fibers with enhanced properties for high-precision sensor and actuator applications.

1. INTRODUCTION Lead zirconate titanate (PZT)-based ceramics are the most widely used piezoelectric until now because of their high piezoelectric response and large scale production capability, and the possibility of tailoring their properties through composition.1 Recently, the European Union has published a health normative (Restriction of Hazardous Substances, RoHS)2 avoiding the use of lead due to its toxicity and environmental risks. Nevertheless, PZT ceramics are temporarily tolerated because of the lack of an adequate alternative. Lately, attention has been moved to lead-free ceramics,3−8 especially potassium sodium niobate (KNN)-based materials. KNN is a good candidate for the replacement of lead zirconate titanate,8−11 but its properties have to be increased in order to reach the PZT ones. Recent efforts contributing to property enhancement in KNN ceramics have been concentrated on chemical optimization through doping,5,8−14 in addition to control over the sintering process,2,3 and domain engineering by the control of the poling process.10,14,15 However, implementations of modified KNN ceramics for commercial use are still limited by their inferior electrical and electromechanical properties as compared to their conventional PZT counterparts. Exceptionally high piezoelectric properties were reported in (K,Na)NbO3−LiTaO3−LiSbO3 (KNL-NTS hereafter), a system proposed by Saito et al. in 2004.9 This pioneering study was based on chemical modifications in the vicinity of the morphotropic phase boundary (MPB) of (K0.5Na0.5)NbO3 by complex simultaneous substitutions in © XXXX American Chemical Society

the A (Li) and B (Ta and Sb) sites of the perovskite lattice. Piezoelectric coefficient, d33, over 400 pC/N was reported for textured ceramics prepared by a complex processing method. More recently, new material systems based on KNN ceramics have been designed and developed for obtaining a new phase boundary and a large d33.16,17 Those ceramics possess a good comprehensive performance of d33, which was obtained due to the design and the optimization of a rhombohedral−tetragonal (R−T) phase boundary. So, these results have opened up a challenge to obtain lead-free piezoceramics with good properties. Indeed, piezoceramic systems are technologically interesting. So, actuator applications require electromechanical coupling providing high strain with high force, e.g, fuel injection, ultrasonic motor, etc. This requirement is largely fulfilled by piezoelectric materials, which allow direct conversion between electrical and mechanical energy. There are a variety of forming methods that have been developed over the years that have been successfully used in compacting the powders to a specific form or shape prior to densification, which is used in obtaining ceramic actuators. Cold pressing in a steel mold is, perhaps, the oldest and most economical of these methods in which ceramic actuators required machining to achieve their final shape. However, the new technological trends evolve toward the integration and Received: April 29, 2014 Revised: May 27, 2014

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dx.doi.org/10.1021/cm501538x | Chem. Mater. XXXX, XXX, XXX−XXX

Chemistry of Materials

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miniaturization of the conventional electronic devices, and thus, more complex procedures in the manufacturing of the actuators are required. In that sense, ceramic fibers and thick patches are required for integrated actuators. Integrated piezoceramics are shaped nearly to their final form and their major drawback is related to deleterious compositional deviations that occur mainly from volatilization during the sintering step. Piezoceramic fibers are used in an ultrasonic transducer, in hydrophones, and in so-called active fiber composites (AFC). In those applications, small fiber diameters are desirable to enable higher achievable resolutions or good flexibility. During the last decades, several techniques have been developed to produce piezoceramic fibers: extrusion (melt spinning),18,19 viscous suspension spinning (VSSP),20 cellulose spinning (ALCERU),21 and sol−gel processing.22 Whereas for PZT fibers only VSSP and ALCERU routes are commercially available, sol−gel fibers have not reached commercial maturity because of straightness and brittleness. Herein, we present a facile method to design high-quality KNL-NTS ceramic fibers with well-regulated K/Na ratios, which is based on a thermoplastic extrusion technique. Two chief achievements were made. One is the clarification of the morphology-inheriting mechanism for obtaining high-quality KNL-NTS ceramic fibers with ideal configuration, while the other is the development of a key sintering route for accurate composition control of K/Na ratios as well as an adequate densification process, which greatly affects the electromechanical functionality. It should be emphasized that an individual KNL-NTS fiber with the optimal densification can exhibit a significant high strain with Smax as high as 0.17% at 3 kV mm−1.

Figure 1. Manufacturing process KNN fibers: (a) photograph of the manufacturing process KNL-NTS fibers. The green fibers are obtained by thermoplastic extrusion as shown in panel a. (b−c) SEM images of a KNL-NTS green fiber where the morphology and the diameter size of the KNL-NTS fiber are shown. The average diameter of the fibers is ∼500 μm. (d) SEM images of the inside of a KNL-NTS green fiber, where 50% of ceramic powders (KNL-NTS nanoparticles) is well distributed with the thermoplastic matrix. The inset shows the agglomerate size distribution (ASD) of KNL-NTS nanoparticles. These KNL-NTS nanoparticles are composed of spherical-like nanoparticles with sizes between 50 and 70 nm,21,22 which form strongly bound agglomerates with sizes of ∼300 nm.

2. RESULTS AND DISCUSSION 2.1. Manufacturing Process, Thermal Behavior, and Morphology of the Green Fibers. Figure 1a shows the manufacturing of KNL-NTS fibers. After passing the dye head of the extruder, the fiber is collected on a conveyor belt. The morphology of the green fibers was characterized by FESEM (Figure 1b−d). The typical geometry of the green fiber in terms of diameter is 500 μm (Figure 1b−c). From Figure 1b, it can be observed that the green fiber possesses a defect-free surface and well-preserved fiber morphology. Figure 1c presents the fracture surface of a green fiber obtained by the thermoplastic extrusion technique. The green fiber is composed of particles corresponding to KNL-NTS calcined powders which are interconnected with the polymeric matrix (Figure 1d). The KNL-NTS nanoparticles have sizes of 50−70 nm, but individual primary particles were clustered into agglomerates, as has been published by Rubio-Marcos et al. 23,24 The agglomerate size distribution (ASD) of the ceramic powder is shown in the inset of Figure 1d. From the ASD, it can be observed that KNL-NTS nanoparticles form agglomerates of ∼300 ± 100 nm (no sintering necks are observed between individual nanoparticles) size distribution. It is surprising that nanopowders are obtained by the solid state reaction process. The origin of such behavior is attributed to the decomposition reaction from carbonates to oxides, which may produce the reduction of their particle size.23,24 A thermal analysis method (TG) was used to establish the optimal debinding procedure. The fiber weight loss reaches ∼16.8 wt % at 900 °C, which correlates with the theoretical weight loss (16 wt %) for the polymeric binder of the green fiber, as shown in Figure 2a. The obtained TG curve (black trace), shown in Figure 2a, suggests only one sharp weight loss

Figure 2. Thermal behavior of the KNL-NTS fibers: (a) TG curve of KNL-NTS green fiber. Only one weight loss can be observed on the TG curve, which is associated with the debinding procedure. Panel b shows the average pore diameter distributions of KNL-NTS nanoparticles, which were thermally treated at 500 °C. From b, it can be observed that the pore size distribution (PSD) is unimodal and is composed by pores with an average diameter of ∼85 nm. The insert of panel b shows a small portion of the porosity located at ∼28 nm. Panel c shows a schematic diagram of the formation process of the porosity after thermal treatment at 500 °C of the ceramic fiber.

stage, which occurs during the progressive heating of the green fiber. The weight loss peak can be observed on the derivative of the TG curve (red trace) at 412 °C, associated with the decomposition temperature of the polyethylene binder. These facts are indicative that the thermal treatment should be B



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performed at temperatures over 412 °C. On the basis of these results, the heat treatment was selected at 500 °C. In this way, we guarantee the complete decomposition of the polymeric binder used in obtaining the KNL-NTS ceramic fibers. Mercury porosimetry measurement was conducted to quantitatively investigate the porosity properties of the ceramic fiber, which was thermally treated at 500 °C. The pore size distribution (PSD) is showed in Figure 2b. The mean pore size locates at ∼85 nm, with size distribution between 55 and 140 nm. This population is associated with porosity between the particle agglomerates, which are formed by the complete decomposition of the polymer precursor during thermal treatment, as mentioned in Figure 1. Moreover, a small content of porosity is located at ∼28 nm, as shown in the inset of Figure 2b. This pore size can be associated with the porosity between the primary particles of the agglomerates. Figure 2c gives a schematic view of the bimodal pore size distribution on the ceramic fiber. Figure 2c-1 and c-2 shows the magnification of the green fiber, which illustrates the polymeric matrix and the agglomerates of ceramic nanopowders of ∼300 nm. These agglomerates are composed of primary particles in the range between 50 to 70 nm,24 assuming that the individual particles have a spherical shape and the same size. When the green fiber is thermally treated at 500 °C, the complete decomposition of the polymeric binder is provoked (Figure 2c), and thus, the porosity formation occurs on the fiber. According to this illustration, the formation of two types of porosity, (i) a large porosity as a result of the removal of polymeric binder and (ii) a small portion of the porosity associated with the space between primary particles with an average pore size of 28 nm, can be explained. 2.2. Correlation between the Structure, Microstructure, and Composition of the Ceramic Fibers. The debound fibers were sintered at temperatures between 1000 and 1200 °C in order to consolidate the ceramic fibers. The Xray diffraction patterns of ceramic fibers at sintering temperatures from 1000 to 1200 °C are shown in Figure 3. The diffraction patterns correspond to a perovskite phase. At lower sintering temperature, the pattern also reveals a minor secondary phase, which was assigned to K 3 LiNb 6 O 17 (PDF#36-0533), with tetragonal tungsten−bronze structure, TTB.25 In a previous study concerning the KNL-NTS system which was conventionally obtained on bulk ceramics, we have demonstrated that the TTB secondary phase possesses a high Li+ content that is incorporated into the perovskite phase when the secondary phase decreased. A relative increase in the sintering time or/and an increase in the sintering temperature resulted in a decrease of the TTB phase content. Therefore, this fact would explain the reduction of the TTB secondary phase on the sintered fibers at temperatures above 1000 °C (Figure 3). The inserts of Figure 3 displays the splitting of the (200) pseudocubic peak into (200) and (002), which suggests a noncubic symmetry in these fibers. The figure shows details of the XRD diffraction pattern in the 2θ range 44.5° to 47° of the KNL-NTS ceramic fibers. In the ceramic fibers sintered at 1000 °C appears a fifth peak that is associated with the occurrence of the secondary phase and is indicated with a triangle symbol. The splitting of the (200) pseudocubic peak is affected by the sintering process. All ceramic fibers show this splitting, which indicates the coexistence between a tetragonal symmetry, T, and an orthorhombic symmetry, O, (see insert of Figure 3). The coexistence of different polymorphs (tetragonal and

Figure 3. Influence of the sintering temperature in the crystalline structure of the KNL-NTS fibers. The figure shows X-ray diffraction patterns of KNL-NTS fibers at sintering temperatures from 1000 to 1200 °C. The inserts of the figure show details of the XRD diffraction pattern in the 2θ range 44.5° to 47° of the KNL-NTS ceramic fibers. These patterns are fitted to the sum of four Lorentzian peaks, which are indexed as 2 tetragonal peaks plus 2 orthorhombic peaks of the perovskite phase. Finally, in the ceramic fibers sintered at 1000 °C appears a fifth peak that is associated with the occurrence of the secondary phase and is indicated with a triangle symbol. (T, tetragonal symmetry; and O, orthorhombic symmetry).

orthorhombic phases) was previously reported on KNL-NTS bulk ceramics.26 It is well-known that the tetragonal symmetry of the perovskite phase can be deconvoluted in two peaks, (002)T and (200)T, fit to a Lorentzian. However, in these patterns two more peaks are also well-defined that appear at ∼45.4 and ∼45.6° (2θ), which are associated with the orthorhombic symmetry. At low sintering temperature (≤1000 °C), the peaks associated with orthorhombic symmetry are more relevant than in the fibers sintered at temperatures between 1100 and 1125 °C, which probably implies a stabilization of orthorhombic symmetry (insert of Figure 3). As a consequence, the most probable origin of this behavior must be related to chemical homogenization with the sintering temperature. As mentioned above, the higher Li+ content of the TTB phase must be incorporated into the perovskite crystalline lattice and thus favored the tetragonal distortion.25,27 So, with the removal of the TTB phase due to the increase in the sintering temperature (in the range from 1100 to 1125 °C), the tetragonal phase is more pronounced. Nevertheless, the system again evolved toward a higher concentration of the orthorhombic phase when the sintering temperature reaches 1200 °C (as shown in the inset in Figure 3). In this case, the origin of such an evolution could be associated with the prolonged exposure of the ceramic fibers at high temperatures during the sintering step (≥1200 °C), which promotes the volatilization of the alkaline elements on the KNL-NTS system.28 The FE-SEM images, shown in Figure 4a−d, illustrate the microstructure of the KNL-NTS ceramic fibers, which were sintered at different temperatures ranging from 1000 to 1200 C



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°C for 2 h, respectively. From Figure 4a.2−d.2, it can be observed that the ceramic fibers have dense microstructure and a typical feature of quadrate shaped grains, which are common in KNN-based ceramics. Grain size distributions, GSD, of the ceramic fibers are shown in the insets of Figures 4a.1−d.1. In the GSDs, it can also be observed that the grain size increases with the sintering temperature. The average grain size (AGS) increases from ∼0.35 ± 0.15 μm in the fibers at a lower sintering temperature (1000 °C) to ∼1.45 ± 0.40 μm for a sintering temperature of 1125 °C. However, the fibers again evolved toward a lower average grain size, ∼1.20 ± 0.20 μm, when the sintering temperature reaches 1200 °C, as shown the inset in Figure 4d.1. Figure 4a.3−d.3 shows details of the morphology of KNL-NTS grains as a function of the sintering temperature. The decrease in grain size at 1200 °C can be explained by the volatilization effect of potassium as will be shown later. The limitation on grain growth is attributed to the existence of grains with very low curvature in spite of their low size. In addition, the straight grain boundaries observed are characteristic of systems in which the sintering occurs through a liquid phase presence.11,29 However, the liquid phase on these samples is hard to detect, implying that it could be a transient liquid phase formed during sintering, with high solubility in the system that led to its eventual disappearance with sintering time and/or sintering temperature.29 In the present study, we have been able to find the presence of a liquid phase located in the grain boundary for the ceramic fibers sintered at 1100 °C (see the yellow arrows in Figure 4b.2 and b.3). The formation of a limited amount of liquid phase during sintering promoted the grain growth by increasing the solution−precipitation mechanism that resulted in cube-shaped grains. The appearance of liquid phases in the grain boundaries at the beginning of the sintering is an experimental proof that the sintering process is mediated by transient liquid phase. Therefore, the purpose of this transient liquid phase is grain growth, which occurs more rapidly in the temperature range between 1000 and 1100 °C. At higher sintering temperatures, this liquid phase is already consumed, and therefore, it cannot be detected. To investigate also the surface morphology of the ceramic fibers as a function of the sintering temperature, a lowresolution scanning electron microscope (LR-SEM) was employed. We note from the LR-SEM images in Figure 1S (Supporting Information) that the ceramic fibers sintered at temperatures ranging from 1000 to 1125 °C possess a defectfree surface and well-preserved fiber morphology (Figures 1Sa− c, Supporting Information). Whereas, the thermal treatment at 1200 °C generates the formation of defects in the fiber surface, as shown in Figure 1Sd (Supporting Information). These defects are provoked by the volatilization of the alkaline elements, which cause an accumulation of gas inside the fiber resulting in the appearance of cracks and protrusions, as indicated with yellow and red arrows, respectively. The main physicochemical characteristics of the ceramic fiber thermal treatment at different temperatures ranging from 700 to 1200 °C during 2 h are presented in the Figure 5a−d. The evolution of the average grain size as a function of the sintering temperature is displayed in Figure 5a. With increasing temperature (see Figure 5a), the average grain size reveals a narrow maximum centered at about 1125 °C and then decreases at 1200 °C. This behavior is therefore consistent with a the volatilization of the alkaline elements at higher temperatures, as ascertained in the XRD data as well with

Figure 4. Microstructural characterization of the KNL-NTS fibers sintered at different temperatures. Microstructure of polished and thermally etched surfaces of KNL-NTS sintered fibers at different temperatures during 2 h: (a) 1000 °C, (b) 1100 °C, (c) 1125 °C, and (d) 1200 °C. Additionally, on the left of each image, the sintering temperature value is indicated. The inserts (a.1 and d.1) show the grain size distributions of KNL-NTS fibers sintered, while panels a.2− d.2 and a.3−d.3 show details of the morphology of KNL-NTS grains. As can be observed from panels a.2−d.2 and a.3−d.3, the samples have a dense microstructure with cuboidal shaped grains, which is a common feature in KNL-NTS based ceramics. In addition, the thermal treatment at 1100 °C generates the partial formation of a liquid phase during the sintering process, which is indicated with yellow arrows in b.2 and b.3. D



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shrinkage for these fibers was calculated to be close to 20%, according to the equation reported in ref 30. In Figure 5b, the dotted line represents the theoretical shrinkage (%). Table 1 Table 1. Effective Shrinkage, Relative Density, and Geometrical Porosity as a Function of Sintering Temperature sintering temperature (°C)

effective shrinkage (%)

relative density (g/cm3)

geometrical porosity (%)

1000 1100 1125 1200

3.9 17.2 18.2 13.5

2.57 4.02 4.17 3.53

44.7 11.9 8.6 22.7

reports effective shrinkage, relative density, and geometrical porosity calculated for every sintering temperature by measuring the fiber diameter before and after heat treatment. Thus, from Figure 5b it is clear that an almost complete densification occurs for the fibers sintered from 1100 and 1125 °C. The porosity also confirms densification of such fibers (see Figure 5b): so, the lower porosity is found in the ceramics sintered at 1100 and 1125 °C. It is known that the K/Na ratio control is critically important for achieving high piezoelectricity in the KNN system,28 but the quantitative measurement of K and Na elements is not easy. However, we have used EDS analysis for the semiquantitative determination of K/Na ratios in resultant KNN-based ceramic fibers. Figure 5c shows the variation of the Na/K ratio as a function of the sintering temperature in the KNL-NTS ceramic fibers. The Na/K ratios shown in Figure 5c are derived from EDS spectra, which represent the atomic percentages of elements. Moreover, the dotted line in Figure 5c represents the theoretical Na/K ratio. Grain compositions of the ceramic fibers are determined by EDS analysis, marked as regions 1, 2, 3, and 4 in Figure 2S (Supporting Information). At thermal treatments in the range between 1000 and 1125 °C, the matrix grains (regions 1, 2 and 3, Figure 2S, Supporting Information) have the typical nominal composition with a Na/K concentration ratio between 1.20 and 1.10, which is close to the nominal ratio of 1.18 (see the table in Figure 2S, Supporting Information). In contrast, for the thermal treatment at 1200 °C according to the EDS analysis (region 4, Figure 2S, Supporting Information), the Na/K concentration ratio (∼2.40) is quite larger than the nominal ratio. Our results reveal that the grains of the fibers sintered at 1200 °C present a sodium-rich composition. This behavior is an experimental evidence of the preferential volatilization of potassium. This dependence (Figure 5c) is in good accordance with plots in Figure 5a and b indicating quantitative similarity in the microstructural behavior and the compositional variety of the ceramic fibers. Figure 5d illustrates average pore diameter, which is derived from the pore size distribution (PSD) of the ceramic fibers depending on the sintering temperature, as shown in the inset to the right of Figure 5c. Indeed, we have to note that the PSD reflects the open porosity in the surface of the fibers. All ceramic fibers, except the sintered fiber at 1000 °C, show a bimodal pore distribution indicating a coexistence of two kinds of porosity populations. So, in Figure 5d we have represented the main population trend of pore diameter with fill symbols, while the open symbols show the evolution of the minority population of the pore diameters. The main population features

Figure 5. Correlation between the microstructural evolution and the sintering process of the KNL-NTS ceramic: (a) evolution of the average grain size as a function of the sintering temperature. The regions labeled ①, ②, and ③ correspond to the different sintering processes of the KNL-NTS ceramic fibers. (b) Effect of the sintering temperature on the radial shrinkage (%) of the KNL-NTS ceramic fibers. The dotted line represents the theoretical shrinkage. (c) Variation of the Na/K ratio as a function of the sintering temperature in the KNL-NTS ceramic fibers. The Na/K ratios shown in panel c are derived from EDS spectra, which represent the atomic percentages of elements (for more information about the composition of the KNLNTS ceramic fibers, see Supporting Information, Figure 2S). In addition, the dotted line represents the theoretical Na/K ratio. (d) Pore diameter evolution with the sintering temperature of the KNLNTS ceramic fibers. In panel d, the filled symbols represent the main population trend of pore diameter, while the open symbols show the evolution of the minority population of the pore diameter in the sample with a bimodal distribution. The minority population of pore diameter at temperatures below 1000 °C is the porosity between the primary particles of the agglomerates. The one at high temperature might come from open porosity at the surface. The color of the pore diameter is correlated with the pore volume and the sintering temperature. Finally, the right inset shows the pore size distribution (PSD) of the KNL-NTS ceramic fibers depending on the sintering temperature. (For more information about the pore volume as a function of the sintering temperature, see Supporting Information, Figure 3S).

microstructure analysis as identified with the formation of defects in the fiber surface and a higher concentration of the orthorhombic phase. The sintering temperature effects on the ceramic fibers are also reflected in the radial shrinkage (%). The theoretical E



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density KNL-NTS ceramic fibers are perfectly adapted to measure the ferroelectric functionality and try to correlate such properties to the observed structural evolution. For a better understanding of the relationship between physical−chemical phenomena and electromechanical properties of the ceramic fibers, the dependence polarization (P − E) as well as bipolar strain hysteresis loops as a function of the sintering temperature were measured, as displayed in Figure 6a−b. In Figure 6 and

evolve with increasing temperature (Figure 5d), the main pore diameter undergoes a drop to values close to ∼9 nm (1125 °C) and then increases gradually culminating in the range of about ∼100 nm, for the fiber sintered at 1200 °C. A similar but less pronounced evolution is observed in the minority population of the pore diameters. Details of the PSD of the KNL-NTS ceramic fibers are shown in Supporting Information, Figure 3S. In comparison to Figure 2, pore size decreases from 500° to 700 °C, which can be explained by the coalescence of the primary particles. At 1000 °C, the porosity inside the agglomerates disappears, and slightly larger porosity occurs because the presintered particle network will not rearrange. This can be confirmed only by small fiber shrinkage of 5% at 1000 °C. At higher temperatures, the minority population might be causing the surface porosity. The increase of the porosity at 1200 °C can be explained by a different sintering mechanism (potassium volotalization). Finally, we have assigned in Figure 5a regions labeled ①, ②, and ③, which correspond to the different sintering processes of the KNL-NTS ceramic fibers. From the perspective of the microstructure and keeping in mind the stages of sintering indicated in Figure 5a, the experimental results observed here can be interpreted as follows. The initial stage, the region marked as ① in Figure 5a, would begin as soon as some degree of atomic mobility is achieved, and during this stage, sharply concave necks begin to form between the particles.31 Comparing the PSD evolution of thermal treatment between 700 and 1000 °C, two remarkable observations can be made. First, the small pores disappear due to the consolidation of the primary particles in the agglomerates; and second, the coalescence of the primary particles allows the formation of grains with sizes of ∼350 nm, provoking an increase in the porosity of the main population. The amount of densification is small, typically the first 5% of the linear shrinkage, which is in agreement with the results obtained in Figure 5b for sintering temperatures between 700 and 1000 °C. In our case, the intermediate stage and final stage occur in thermal treatments between 1000 and 1125 °C (region ② of Figure 5a). The intermediate stage is characterized by moderating the high curvature of the initial stage, and the microstructure consists of a three-dimensional interpenetrating network of particles and continuous channel-like pores. Grain growth (coarsening) starts to become significant. Besides, at these temperatures a transient liquid phase appears, as observed in Figure 4b, which accelerates the coalescence and grain growth processes. Thus, an evidence of this behavior is found on the increased average grain size at temperatures between 1000 and 1100 °C, where the average grain size increases from ∼ 0.35 ± 0.15 μm to ∼1.30 ± 0.30 μm, representing 370%. As sintering proceeds, the channel-like pores break down into isolated, closed voids, which marks the start of the final stage. In this stage, grain growth is more extensive, and difficulties are commonly encountered in the removal of the last few percentages of porosity. This takes place in our fibers for thermal treatment between 1100 and 1125 °C. Eventually, this system can evolve to a last stage that we have denominated region ③ in Figure 5a, which occurs for thermal treatment at 1200 °C. This stage is defined by the volatilization of potassium in the perovskite structure, causing a worsening of the microstructure (i.e., grain size reduction, increased porosity, and appearance of defects on the surface of the fibers). 2.3. Influence of the Sintering Temperature in the Functional Properties of the Ceramic Fibers. The high

Figure 6. Evaluation of the ferroelectric properties of KNL-NTS ceramic fibers: (a) average polarization and (b) strain evolution as a function of the applied field.

Table 2, the electromechanical properties of the single KNLNTS fibers are reported. As expected, for sintering temperatures corresponding to 1000 and 1200 °C low properties are achieved. Almost no remnant polarization and piezoelectric coefficient could be measured. For intermediate temperatures, rectangular polarization curves are obtained with relatively low coercive fields. In these conditions, maximum strains (Smax) as high as 0.17% could also be obtained (at 1100 °C). The strain behavior is a key feature of actuator systems. Nevertheless, this behavior is only occasionally studied. Up to now, typical strain values for KNN-based ceramics are around 0.13%32 and 0.16%33 under electric fields of 4 and 6 kV mm−1, respectively. Unfortunately, practical implementations of KNN-based ceramics for commercial use are still limited by their inferior electromechanical properties as compared to those of their conventional PZT counterparts.34 In the present study, it is worth pointing out that the strain reaches 0.17% at 3 kV mm−1, which is a very promising value for lead-free piezoceramic fibers. Fibers sintered at 1100 and 1125 °C also show a significant increase in d33: values close to 140 pm/V are registered for the 1125 °C fibers. In bulk KNL-NTS samples with the same composition as the one here reported, an increase in d33 is observed by increasing the sintering time at 1125 °C.25,35 A d33 value close to 140 pm/V was obtained after 8 h of sintering at 1125 °C. The high d33 of 240 pm/V was achieved for bulk material, which was sintered at 1125 °C for 16 h.35 Further research on sintering behavior on KNL-NTS fibers is necessary F



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Table 2. Main Electromechanical Properties of the Fibers Sintered at Different Temperatures sintering temperature (°C)

Pr (μC/cm2)

1000 1100 1125 1200

0.6 ± 0.3 13.3 ± 2.2 20.4 ± 1.1 2.7 ± 0.9

Ec (kV/mm) 0.8 1.3 1.2 1.0

± ± ± ±

0.3 0.1 0.1 0.1

Smax (%)

d33 (pm/V)

ε (−)

± ± ± ±

6±2 130 ± 36 139 ± 23 21 ± 6

385 ± 13 1085 ± 75 1455 ± 115 915 ± 4

0.02 0.17 0.13 0.05

0.01 0.02 0.01 0.01

Figure 7. Confocal Raman spectroscopy characterization of KNL-NTS ceramic fibers sintered at 1100 °C: (a) optical micrographs from inside to the edge of the polished ceramic fiber, which was sintered at 1100 °C. (b) The Raman image of the KNL-NTS ceramic fibers exhibiting the appearance of a secondary phase. The Raman map was derived by summing the total spectral pixel intensity from 100 to 950 cm−1. The secondary phase associated with the KNL-NTS phase is indicated in green (*), whereas the KNL-NTS perovskite phase corresponds to red regions. Average Raman spectra of KNL-NTS ceramic fiber sintered at 1100 °C (c.1) and of the secondary phase (c.2). The inserts of c.1 and c.2 show magnified Raman spectra and Lorentzian fits in the frequency range between 450 and 750 cm−1. These spectra are fitted to the sum of three Lorentzian peaks, ascribed to Eg (υ2), A1g (υ1) corresponding to Raman modes of the KNL-NTS perovskite phase and the third peak that is associated with the occurrence of the secondary phase. In addition, the reader can find more information about this secondary phase in Supporting Information, Figure S4.

to be able to achieve comparable d33 coefficients in comparison to bulk samples. It can be observed from Figure 6a that the Pr values depend on the sintering temperature. The Pr value of the fiber sintered at 1100 °C is 13.3 μC/cm2, whereas this value increases to 20.4 μC/cm2 for the fiber sintered at 1125 °C. In addition, a reduction of the coercive field occurs for the fiber sintered at 1125 °C. From this result, slightly higher Smax in the fiber thermal treatment at 1125 °C is expected, contrary to the observed values (see Table 2). So, this intriguing improvement of the Smax (for the fiber sintered at 1100 °C) cannot be explained only on the basis of intrinsic effects, considering that the piezoelectric properties of these materials are mainly modulated by the softening of the ferroelectric properties. Therefore, the improvement of Smax should be mainly attributed to extrinsic effects. The main extrinsic effect as the grain size also fails to justify such behavior. The slightly lower porosity could contribute greatly in the high strain by reducing the ferroelastic stress that occurs under the electric field. Moreover, the larger coercive field implies that the domain structure required greater energy to produce domain reversal. Under an applied electric field, the strain in piezoelectric ceramic is mainly attributed to the switching of the 90° domain. The coexistence of tetragonal and orthorhombic symmetries implies the presence of more options to produce strain, based on the appearance of orthorhombic 120 and 60° domains.15,36

Orthorhombic domain growth in previously formed tetragonal domains during the sintering process, the higher density, limited such domain development, as was established by different research groups in KNN-based ceramics.10,15,37 2.4. Characterization of the Stress in Ceramic Fibers by CRM. To verify that the electromechanical properties are influenced by the stress accumulation into ceramic fibers, additional experiments were performed by confocal Raman microscopy (CRM). The fiber thermal treated at 1100 °C was chosen to set the optimal intrinsic and extrinsic conditions at which the strain values were maximized. Figure 7a shows an optical micrograph from inside to the edge of the polished fiber, which was aligned perpendicular to the Raman laser. The 70 × 70 μm area denotes the selected region where the Raman spectra were collected at a plane located just below the surface of the sample, where the Raman intensity was maximized. The acquisition time for a single Raman spectrum was 600 ms; thus, the acquisition of a Raman image consisting of 70 × 70 pixels (4900 spectra) required 49 min. Features such as Raman peak intensity, peak width, and Raman shift from the recorded Raman spectra were fitted with algorithms to compare information and to represent the derived Raman image (Figure 7b). The vibrations of the BO6 octahedron consist of 1A1g (υ1) + 1Eg (υ2) + 2F1u (υ2, υ4) + F2g (υ5) + F2u (υ6) modes. Of these vibrations, 1A1g (υ1) + 1Eg (υ2) + 1F1u (υ3) are stretching modes and the rest bending modes.27 In particular, A1g (υ1) G



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symmetrical mode and F2g (υ5) antisymmetric mode are detected as relatively strong scattering signals in KNN-based materials because of a near-perfect equilateral octahedral symmetry. Raman spectra having the same Raman shift are classified by colors correlating the color intensity with the Raman intensity. A careful examination of the Raman spectra (Figure 7c) reveals the presence of Raman modes corresponding to the KNL-NTS phase (Figure 7c.1) and the appearance of a new peak, as shown in Figure 7c.2. The presence of the main vibration modes of the NbO6 octahedron as shown in Figure 7c.1 is another evidence of the formation of perovskite structure, which corresponds to the red regions marked in Figure 7b. The third peak, indicated as the green region in Figure 7b, is attributed to the appearance of a characteristic Raman mode of the TTB phase at ∼690 cm−1, which emerges concomitantly. The average Raman spectrum of the secondary phase is shown in Figure 7c.2, which can be indexed on the basis of a phase mixture constituted by a majority of KNL-NTS phase and a minority of secondary phase. As alluded to earlier, this secondary phase was not observed by a corresponding Xray diffraction pattern. This could imply that the secondary phase is well below the XRD sensitivity by low amount or low crystallinity. The most likely origin of the secondary phase appearance can be associated with the crystallization of the liquid phase transient, which was observed by FE-SEM (Figure 4b.2 and b.3). Details of this region are presented in the inset of Figure 7c.2, where the spectrum is fitted to the sum of three Lorentzian functions centered at ∼555 cm−1, ∼612 cm−1, and ∼690 cm−1, and ascribed to the Eg (υ2) and A1g (υ1) Raman modes of the KNL-NTS phase and TTB phase, respectively. From the inset of Figure 7c.2, the secondary phase content can be calculated, which represents ca. 1.2% of the total measured phases. Thus, the identification and quantification of minority secondary phases can be approached by CRM. It is worth pointing out other characteristics revealed by the CRM imaging since it has been possible to locate the regions in which the secondary phase coexists with the KNL-NTS phase (marked as green regions in Figure 7b). (More information about this secondary phase can be found in Supporting Information, Figure 4S). A correlation between Raman peak shifts and changes of material stress is established.38 Considering that the promising improvement of the Smax found in this study (0.17% at 3 kV mm−1) is supported by the stress accumulation within the ceramic fiber (i.e., coexistence of tetragonal and orthorhombic symmetries as well as by the inherent fiber morphology), and keeping in mind that Raman spectroscopy is sensitive to stress in ceramic materials, we used confocal Raman mapping to monitor the distortion of the oxygen octahedra originated on the ceramic fiber sintered at 1100 °C. The evolution of the A1g mode shown in Figure 8a was measured within the spatial resolution of the technique following the black arrow in Figure 8b. Clearly, a continuous increase in the Raman shift of this mode occurs from the inside (point marked as A in Figure 8b) to the edge (point indicated as B in Figure 8b) of the fiber sintered at 1100 °C. Figure 8b shows the Raman shift image of the A1g Raman mode for the selected area of the ceramic (Figure 7a). The Raman imaging allows one to evaluate the stress because each pixel comprises a full Raman spectrum. This Raman map shows three featured regions, which were named ①, ②, and ③. These three regions correspond to relaxation zone ①, an intermediate zone ②, and stress zone ③ of the ceramic

Figure 8. Evaluation of stress on a KNN ceramic fiber sintered at 1100 °C by confocal Raman spectroscopy: (a) evolution of the Raman shift of the A1g mode, which were measured following the black arrow in panel b. The black arrow represents the distance from the center, point A, to the fiber surface, point B. (b) Color coded map showing the Raman shift corresponding to the A1g mode. The points labeled 1, 2, and 3 correspond to the different regions shown in panel a. The bar with the color code corresponds to values of the Raman shift determined at each pixel. (c) Magnified Raman spectra and Lorentzian fits of KNN-based ceramic fiber in the 450−750 cm−1 frequency range corresponding to the points labeled 1, 2, and 3 in the images in a and b. These spectra are fitted to the sum of two Lorentzian peaks, ascribed to the Eg (ν2) and A1g (ν1) Raman modes. The insert at the top of panel c shows an enlargement of the Raman shift taken between points 1 and 3, exhibiting an increased Raman shift of 2.5 cm−1.

fiber. The Raman spectra of each region are shown in Figure 8c. From Figure 8b and c, it is worth noticing a characteristic aspect in the A1g Raman shift evolution along the line crossing the fiber. Figure 8b reveals a stress zone (③) that is located on the outside of the fiber which extends toward the inside about ∼15 μm, while the relaxation zone occupies the remaining space of the fiber (in term of diameter is ∼400 μm). This fact suggests that the sintering process takes place from outside to inside the fiber, resulting in a stress gradient within the material. Thus, the relaxation zone is characterized by a suitable AGS (∼1.30 ± 0.30 μm), an appropriate PSD with a pore diameter of ∼10 nm, a convenient composition control, and a small amount of the secondary phase. All these extrinsic factors are responsible for the electromechanical coefficient, Smax, reaching 0.17% at 3 kV mm−1. To date, this value can be considered to be one of the best results reported in KNN-based piezoceramics. Moreover, the high strain obtained is particularly remarkable because it indicates that lead-free-based materials could be processed as low dimensional piezoceramics in a more effective way than classical PZT fibers and thick film materials in which the lead volatilization is environmentally undesirable. The phenomenon described here could be explained by a simple model based on the liquid-phase sintering of ceramic fibers that evolves toward the tuning of the microstructure during the thermal treatment. The tuning of the microstructure is the reason to demonstrate a high level of unipolar strain up to 0.17% at 3 kV mm−1. So, Figure 9 shows a schematic representation of the sintering stages. At temperatures between 500 and 700 °C, the debinding procedure occurs. This process H



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Figure 9. Schematic representation of the sintering mechanism on the KNL-NTS ceramic fibers. The scheme illustrates the sintering process of the ceramic fibers, which are characterized by three steps: (1) initial stage, (2) intermediate stage, and (3) last stage. (A detailed explanation of the scheme can be followed in the main text.)

properties of fibers sintered at 1100 and 1125 °C is the relationship to internal stress, which is determined by confocal Raman spectroscopy. According to the results, it could be investigated that internal stress on fibers sintered at 1125 °C is higher. This internal stress can provoke a blocking of the strain at higher voltages, and therefore, lower longitudinal free strain occurs. A schematic representation of the sintering mechanism of the KNL-NTS ceramic fibers is also established. So, the control of the sintering process allows for correlating the extrinsic factors of the system with the high value of Smax, which is found for the KNL-NTS ceramic fibers sintered at 1100 °C. We believe that the general strategy and design principles described in this study will open new avenues in developing of (K,Na)NbO3based lead-free piezoelectric fibers with enhanced properties for high-precision sensor and actuator applications.

is characterized by the decomposition of the polymeric binder. At these temperatures, the ceramic fibers are composed of spherical-like nanoparticles with sizes between 50 and 70 nm, which form strongly bound agglomerates with sizes of ∼300 nm. Increasing the sintering temperature to 1000 °C, the agglomerates consolidate and lead to the formation of grains with sizes of ∼300 nm (denoted as ①). The amount of densification is small, typically the first 5% of the linear shrinkage which is in agreement with the results obtained in Figure 5b for sintering temperatures between 700 and 1000 °C. At 1100 °C, grain growth (coarsening) starts to become significant. Besides, at these temperatures the transient liquid phase observed in Figure 4b appears, which accelerates the coalescence and grain growth processes. So, this stage is defined by disordered matrix grains with an average size of around 1.3 μm, including the appearance of the transient liquid phase and the formation of large pores, as shown in Figure 4 (this stage is denominated as ②). With increasing temperature to 1125 °C, grain growth is more extensive, and difficulties are commonly encountered in the removal of the last few percentages of porosity (minimal porosity). At the optimum sintering temperature (1125 °C), the system reaches its maximum grain growth with values ∼1.4 μm. Finally, increasing the sintering temperature to 1200 °C causes decomposition of the system (stage labeled as ③), due to the volatilization of alkali elements (see Figure 4d). This might be the main reason why lower densities in this ceramic fiber were measured, producing a decrease in average grain size to ∼1.2 μm. Under these sintering conditions, besides the above extrinsic factors, KNL-NTS ceramic fiber is characterized by a convenient composition control, as well as a surprising stress gradient within the material, as shown the Raman analysis. This appropriate configuration of extrinsic factors is responsible for Smax reaching one of the highest values reported in KNN-based lead-free piezoelectric fibers. Thus, we can modulate the extrinsic factors controlling the sintering process. It is worth pointing out that this stimulant behavior observed here could have potential technological applications leading to the understanding of some key features for actuator systems.

4. EXPERIMENTAL PROCEDURES 4.1. Sample Preparation. The compositional ceramics of (K0.44Na0.52Li0.04) (Nb0.86Ta0.10Sb0.04)O3, hereafter abbreviated as KNL-NTS, was prepared by conventional solid-state reaction from an adequate mixture of corresponding oxides and carbonates. Na2CO3, Li2CO3 (PANREAC, >99.5%), K2CO3 (Merck, >99%), Nb2O5, Ta2O5, and Sb2O5 (SIGMA-Aldrich, > 99.9%, >99% and >99.995%, respectively) were used as starting raw materials. In all of the experiments, the raw materials were dried right before used at 200 °C for 1 h because of their hygroscopic nature.39 Raw materials were milled individually (attrition milling in ethanol for 3 h) in order to obtain an appropriate particle size distribution.24 These powders were then weighed and mixed by attrition milling using ZrO2 balls in ethanol medium for 3 h, dried, and calcined at 700 °C for 2 h with a heating rate of 3 °C/min. 4.2. Fiber Preparation. A 50 vol % feedstock was prepared by mixing twice KNL-NTS powder, polymer (LDPE, PEBD 1700 MN, Lacqtene Elf Atochem S.A., Switzerland), and surfactant (Stearic acid >95%, Sigma-Aldrich) in a high-shear mixer (Haake PolyLab Mixer Reomix 600, Thermo Fisher Scientific, Germany). Green fibers with a diameter of ∼500 μm were obtained by thermoplastic extrusion of the feedstock in a capillary rheometer (RH7 Flowmaster, Malvern, Germany). The pressure during extrusion was recorded by a pressure sensor positioned before the dye entry. The fibers were extruded at a piston speed of 0.814 mm/min and collected on a conveyor belt, where they were cut in ∼10 cm long pieces. After debinding at 500 °C (1 °C/min, 1 h), green fibers were densified at different temperatures: 1000, 1100, 1125, and 1200 °C, keeping the heating rate constant at 3 °C/min. In order to reduce alkali evaporation, the fibers were closed in alumina crucibles, sealed with alumina powder. 4.3. Thermal Characterization. Thermogravimetric analysis was carried out on KNL-NTS ceramic fibers before heat treatment using a NETZSCH STA 409/C analyzer (Netzsch, Germany). The fibers were placed in a Pt/Rh crucible and heated up to 900 °C with a

3. CONCLUSIONS On the basis of KNL-NTS nanopowder, lead-free piezoelectric fibers have been developed. The influence of the sintering temperature on functional properties of the ceramics has been stated. Maximum strain as high as 0.17% is recorded for fibers sintered at 1100 °C, while d33 values close to 140 pm/V are shown by the ones heat treated at 1125 °C. One possibility to explain this antagonism between large signal and small signal I



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heating rate of 3 °C/min. The measurements were performed in a flowing air atmosphere. 4.4. Structural Characterization. The crystalline structure was determined by X-ray diffraction analysis (XRD, X’Pert PRO Theta/ 2theta of Panalytical, Cu Kα radiation, PANalytical, The Netherlands) on powder obtained by milling of the sintered ceramics. The lattice parameters were refined by a global simulation of the full diagram (pattern matching, Fullprof program).40 4.5. Microstructural Characterization. The particle size and morphology of the powder ceramics forming the fiber were evaluated using a Field Emission Scanning Electron Microscope, FE-SEM (Hitachi S-4700, Tokio, Japan), fitted out with energy dispersive spectroscopy (EDS). Sintered and fresh fracture surfaces of the fibers were evaluated. The average grain size and grain size distributions were determined from FE-SEM micrographs by an image processing and analysis program (Leica Qwin, Leica Microsystems Ltd., Cambridge, England) considering more than 250 grains in each measurement. The PSD and volume data were determined using PoreMaster GT, (Quantachrome Instruments, United Kingdom). Thus, pore size is calculated using the Washburn equation: pore diameter = −(4γ cos θ)/P, where γ = surface tension, P = pressure, and θ = contact angle. 4.6. Electrical Characterization. The longitudinal free strain of a single fiber is measured by the use of a special equipment (FerroFib) developed from a collaboration between Empa and aixACCT System GmbH (Germany). A sketch of this instrument is reported in Figure 5S, Supporting Information. A 2.5 mm long fiber is glued vertically on a PMMA disc with silver paste. This glue allows the electrical contact between the front end of the fiber and the metallic bottom electrode. The same silver paste is used to create the electrode on the upper side of the fiber. Top and bottom electrodes are then connected with a probe and needle as shown in Figure 5S, Supporting Information. After the first polarization, when an alternating electric field is applied, the fiber elongates and contracts longitudinally, and the strain/ contraction is recorded by a laser positioned on the top of the probe. The strain (%) and the specific electrical field (kV/mm) can then be calculated by knowing the length of the fiber. For P-E loops and butterfly loops, frequencies of 0.1 and 1 Hz were used by applying a voltage of 3 kV/mm. Smax was calculated as the difference between the minimum and maximum displacement in the positive side of the butterfly curve. Piezoelectric constant d33 was determined at 10 Hz under an applied voltage of ±100 V. It is worthwhile to mention that d33 is frequency dependent.41 For the fibers, only a frequency between 1 and 100 Hz can be measured by the new FerroFib device. 4.7. Characterization of Ceramic Fibers by CRM. The experiments were performed using a Confocal Raman Microscope (Witec alpha-300R, WITec GmbH; Ulm, Germany). Ceramic samples for the CRM were only polished, and no further thermal or chemical etching was used to reveal the domain structure and thus to avoid topographic artifacts due to nanoroughness. With this aim, the surface of the fibers was carefully polished to obtain mirror finish surfaces. As a consequence, the microroughness is inhibited resulting in an improvement of the finish of the sample surface. Raman spectra were obtained using a 532 nm excitation laser and a 100× objective lens (NA = 0.9). The incident laser power was 0.5 mW. The optical diffraction resolution of the confocal microscope was limited to about ∼200 nm laterally and ∼500 nm vertically. Raman spectral resolution of the system was down to 0.02 cm−1. The microscopy sample was mounted on a piece-driven scan platform having 4 nm lateral and 0.5 nm vertical positional accuracy. The microscope base was also equipped with an active vibration isolation system, active at 0.7−1000 Hz. Collected spectra were analyzed by using Witec Control Plus Software.

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pores in the sintering process of the KNL-NTS ceramic fiber; characterization of the secondary phase observed in the ceramic fiber sintered at 1100 °C through confocal Raman spectroscopy; and sketch of the FerroFib equipment. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*Phone: +34 91 735 58 40. Fax: +34 91 735 58 43. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We acknowledge the COST Action MP0904 for financing the STSM and the Marie Curie Cofund project (EmpaPostDoc program) for general financial support. We also express our thanks to the MICINN (Spain) project MAT2010-21088-C0301 for their financial support. F.R.-M. is also indebted to CSIC for a ‘‘Junta de Ampliación de Estudios’’ contract (ref JAEDOC071), which is cofinanced with FEDER funds. We thank C. A. Fernández-Godino for his assistance in the preparation of the compositions.

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

Microstructural evolution of the ceramic fiber surface as a function of sintering temperature; composition evaluation in the KNL-NTS fibers sintered at temperatures different from EDS analysis; evolution of the volume and the diameter of the J



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High Strain in (K,Na)NbO3-Based Lead-Free Piezoelectric Fibers

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