Morphology of Porous Hosts Directs Preferred Polymorph Formation

Article pubs.acs.org/crystal

Morphology of Porous Hosts Directs Preferred Polymorph Formation and Influences Kinetics of Solid/Solid Transitions of Confined Pharmaceuticals Gitte Graubner,† Gopalakrishnan Trichy Rengarajan,† Nicole Anders,‡ Nicole Sonnenberger,§ Dirk Enke,‡ Mario Beiner,§,# and Martin Steinhart*,† †

Institut für Chemie neuer Materialien, Universität Osnabrück, Barbarastrasse 7, 49076 Osnabrück, Germany Universität Leipzig, Institut für Technische Chemie, Linnestrasse 3-4, 04103 Leipzig, Germany § Martin-Luther-Universität Halle-Wittenberg, Naturwissenschaftliche Fakultät II, 06099 Halle (Saale), Germany # Fraunhofer IWM, Walter-Hülse-Strasse 1, 06120 Halle (Saale), Germany ‡

S Supporting Information *

ABSTRACT: The pore morphology of a porous host may determine which polymorph a crystallizable guest preferentially forms and may influence the kinetics of solid/solid transitions. Slow cooling of the drug acetaminophen (ACE) inside the straight cylindrical pores of anodic aluminum oxide (AAO, tortuosity = 1) in contact with a bulk ACE surface film preferentially yields uniformly oriented form II and/or form III crystals. The occurring orientations of form II and form III crystals are characterized by high structural registry along the AAO pores. The uniformly oriented form III crystals inside the AAO pores were readily converted into likewise uniformly oriented form II crystals by a solid/solid transition. Thus, we obtained uniformly oriented form II crystals in AAO at high yields. We suggest that sporadic heterogeneous nucleation at bulk crystals formed in the bulk surface film on top of the AAO coupled with kinetic selection of crystal orientations results in fast growth of properly oriented crystals along the 100 μm deep AAO pores. This mechanism is suppressed in controlled porous glass (CPG) having isotropic spongelike pores (tortuosity > 1.5) with free growth paths on the order of 100 nm, where form I formed instead. Moreover, the transition from form III to form II is suppressed in CPG. Possible reasons may include impingement of the propagation front of the solid/solid transition on the CPG pore walls after short propagation paths and inevitable formation of form II grains with different orientations separated by energetically disadvantageous grain boundaries. The results reported here are relevant to mesoscopic crystal engineering aimed at controlled drug release from nanoscale delivery systems. Polymorphs not accessible otherwise in nanoscale containers may be produced at high yields. The principles reported here may be transferred to areas such as nanowire-based organic electronics.

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morphologies and that pore morphology may affect solid/solid transitions between different polymorphs. Melt crystallization and cold crystallization of polymorphic materials confined to straight cylindrical pores of a porous host may occur in the presence of a bulk surface film of the crystallizing material on top of the porous host. Inside straight cylindrical pores filled with supercooled or vitrified melt, new crystals may sporadically form close to the pore mouths by heterogeneous nucleation on interfaces to bulk surface film crystals impinging on the porous hosts. The thus-emerging crystals may have different orientations and may even belong to other polymorphs than the bulk surface film crystals. Any crystal slowly growing along straight cylindrical pores may in turn act as a nucleation site for

onfinement imposed by nanoscale drug carriers influences crystallization of polymorphic drugs.1 However, predictive understanding as to how drugs crystallize under confinement, although of tremendous importance for controlling their solubility, release kinetics, and bioavailability, has remained premature.2 Solubility studies of drugs are often elusive because they are not polymorph-specific. Moreover, the morphological importance of specific crystal faces may affect release kinetics in the case of nanoscale drug crystals with large surface-to-volume ratios. Rodlike nanoscale drug carriers show several advantages, such as extended residence times3 and superior internalization rates.4 Improved understanding of crystallization of polymorphic drugs in specific confining geometries is, therefore, a prerequisite for rational drug release management from drug carrier systems combining the advantages of nanoscopic size, anisotropic shape, and mesoscopic crystal engineering. Here we show that polymorphic drugs may preferentially form different polymorphs in porous hosts with different pore © 2013 American Chemical Society

Received: August 1, 2013 Revised: November 25, 2013 Published: November 25, 2013 78

dx.doi.org/10.1021/cg401170m | Cryst. Growth Des. 2014, 14, 78−86

Crystal Growth & Design

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dominance of fast-growing crystals having a specific orientation with respect to the pore axes is suppressed. Different polymorphs than in straight cylindrical pores may prevail in spongelike pores even if crystallization is otherwise carried out under the same conditions. Likewise, solid/solid transitions in properly oriented crystals proceeding along cylindrical pores may be suppressed in spongelike pores by restricted propagation paths and possibly by formation of energy-rich boundaries between grains of the newly formed polymorph with different orientations. We used the antipyretic and analgesic drug acetaminophen (ACE) as model system. We compared crystallization of ACE in straight and aligned cylindrical pores of self-ordered anodic aluminum oxide (AAO)5 (Figure 2a) and in isotropic, continuous (spongelike) pore systems of controlled porous glass (CPG)6 (Figure 2b). Both host systems had pore diameters of 60 nm. The free growth path in CPG is a few tens of nanometers. The free growth path in AAO pores of ∼100 μm corresponding to the AAO pore depth is 3 orders of magnitude longer than in CPG. Both template systems have oxidic pore walls. ACE forms rigid amorphous interphases on the pore walls of silica hosts.7 As evidenced by differential scanning calorimetry (DSC), significant crystallization of ACE confined to AAO pores without contact to bulk ACE surface films on top of the AAO did not occur, independent of the applied thermal treatment.8 Therefore, we conclude that in AAO and CPG ACE/pore wall interactions are similar and characterized by the presence of a rigid amorphous ACE layer, while surface nucleation of ACE on the pore walls can be ruled out.

further heterogeneous nucleation events. Crystal growth rates along directions normal to crystal faces can significantly differ between specific crystal faces of the different polymorphs, and there might be distinct directions of outstandingly fast crystal growth. Occasionally, heterogeneous nucleation will yield crystals having a direction of outstandingly fast crystal growth aligned with the long axes of the straight cylindrical pores (Figure 1a). Once such crystals have formed, they rapidly grow

Figure 1. Crystallization of polymorphic materials in the straight cylindrical pores of anodic aluminum oxide (AAO) and the spongelike pores of controlled porous glass (CPG). AAO and CPG pore walls are black; supercooled or vitrified liquid is light gray; crystals having a direction of fast crystal growth oriented along the long axes of the pores are dark yellow; bulk surface film crystals offering interfaces for heterogeneous nucleation and other crystals are red. (a) New crystals form by heterogeneous nucleation at the interface of bulk surface film crystals having different orientation or consisting even of a different polymorph. If a direction of fast crystal growth (indicated by an arrow) is aligned with the host’s pore axes, the new crystals will rapidly grow along the host’s pores. (b) In AAO (tortuosity = 1) fast-growing new crystals grow along the straight cylindrical pores (free growth path tens of micrometers) until they fill the entire pore volume. (c) In CPG (tortuosity > 1.5) fast-growing crystals rapidly impinge on the walls of the tortuous CPG pores (free growth path tens of nanometers) so that other crystallization mechanisms may dominate.

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METHODS

Materials. Acetaminophen (ACE, C8H9NO2 CAS 103-90-2 purity 99.0%) purchased from Sigma-Aldrich was used without further purification. Self-ordered AAO was prepared according to procedures described elsewhere.5 The cross-sectional scanning electron microscopy (SEM) image of an AAO membrane with a pore diameter of 60 nm shown in Figure 2a was acquired on a JEOL 7500F microscope operated at 2 kV with a below-the-lens secondary electron detector. Controlled porous glass (CPG) with an average pore diameter of 60 nm was prepared by spinodal decomposition of sodium borosilicate initial glasses (composition: 70 wt % SiO2, 23 wt % B2O3, and 7 wt % Na2O) followed by leaching of the formed alkali-rich borate phase. The texture properties of

along the straight cylindrical pores until they occupy the entire pore volumes. Therefore, arrays of aligned straight and cylindrical pores will be filled with uniformly oriented crystals of a specific polymorph on a macroscopic scale (Figure 1b). If crystallization is carried out under the same conditions but in continuous and isotropic spongelike pore systems, fast-growing crystals have only short free growth paths before they impinge on the pore walls (Figure 1c). Hence, the growth mechanism dominating in arrays of straight cylindrical pores and the

Figure 2. Scanning electron microscopy images of rigid inorganic porous hosts with a pore diameter of 60 nm. (a) Cross-sectional view of selfordered AAO; (b) CPG monolith. Both panels have the same magnification. 79



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ACE surface films. Some samples were annealed at 135 °C for 1 h under argon after removal of the bulk ACE surface reservoirs. To produce ACE form II/form III inside AAO by nonisothermal crystallization, ACE was heated to 175 °C for 30 min under argon on top of AAO membranes and cooled to room temperature at a cooling rate of −0.5 K/min. Then, the bulk ACE surface films were removed from the AAO surfaces using sharp blades. To produce amorphous ACE in AAO for isothermal cold crystallization, ACE was infiltrated into AAO at 180 °C. The AAO containing ACE and covered with a bulk ACE surface film heated to 180 °C was rapidly put in a freezer cooled to below −10 °C. To produce form I inside AAO, ACE was infiltrated at 175 °C for 30 min under argon. The ACEinfiltrated AAO covered with a bulk ACE surface film was quickly removed from the furnace and quenched to below 0 °C by putting the sample on cooled copper plates under ambient conditions. Differental Scanning Calorimetry (DSC). Prior to any DSC measurement, bulk ACE surface films were removed with sharp blades. AAO membranes infiltrated with ACE were initially connected with underlying aluminum substrates. The aluminum substrates were etched with solutions containing 1.7 mg of CuCl2·2H2O, 50 mL of deionized H2O, and 50 mL of concentrated HCl(aq) under cooling with ice water. For this purpose, the samples were mounted on specifically designed sample holders that prevented the AAO pore openings and the infiltrated ACE from coming into contact with the etching solution. After the etching step, the samples were dried under a vacuum at room temperature. DSC measurements were carried out with a power-compensation DSC 8500 (Perkin-Elmer) equipped with an Intracooler 3 unit. All heating and cooling scans were performed at rates of ±10 K/min. Wide-Angle X-ray Scattering (WAXS). WAXS patterns were measured in reflection using a PANalytical X’Pert Pro MRD diffractometer operated with Cu Kα radiation at 40 kV and 40 mA. Figure S1, Supporting Information, displays the geometry used to investigate AAO samples. Θ/2Θ scans allowed selective detection of scattering intensity originating from sets of lattice planes oriented parallel to the AAO surface and perpendicular to the AAO pore axes. In the course of a Θ/ 2Θ scan, the AAO was tilted about the Θ axis indicated by the dotted line in Figure S1 by an angle Θ (the scattering angle). Incident beam and AAO surface enclosed the scattering angle Θ. The Θ axis lay in the plane of the AAO surface and was oriented perpendicularly to the AAO nanopore axes and to the scattering plane (defined by incident beam and detector). The detector was rotated about the Θ axis by an angle 2Θ. The obtained Θ/2Θ patterns were background-corrected and analyzed using the program PANalytical X’Pert HighScore. Angles between sets of lattice planes were calculated with the program Diamond 3.0d (Crystal Impact, 2005) based on crystal structures deposited in the Cambridge Structural Database. For Schulz-Scans,11 the diffractometer was configured as follows. The incident beam passed a nickel filter, a polycapillary (length 70 mm), and a cross slit collimator (slit width and height 500 μm). The diffracted beam passed a parallel plate collimator before it was collected with a PW3011/20 proportional point detector. Schulz scans were measured with fixed Θ and 2Θ angles by tilting ACE-infiltrated AAO membranes about the Ψ axis (indicated by a solid line in Figure S1). The Ψ axis lay in the plane of the AAO surface and in the scattering plane but was oriented perpendicularly to the Θ axis and to the AAO nanopore axes. Schulz scans yielded

CPG were determined by mercury intrusion. The measurements were carried out on a Quantachrome PoreMaster. A contact angle of 141.3° for Hg was used for further calculations. The average pore diameter was calculated by applying the Washburn equation and a cylindrical pore model. The secondary electron SEM image of a CPG membrane seen in Figure 2b was obtained with a Zeiss Auriga microscope at an acceleration voltage of 1.5 kV. Tortuosity is the ratio of the mean effective path length between two points in a nanoporous host and the shortest possible distance in the absence of obstacles. Whereas the tortuosity of straight cylindrical AAO pores amounts to 1, the tortuosity of the isotropic, continuous pore systems of CPG was larger than 1.5. Shlekhin et al.9 investigated the tortuosity of porous glass membranes and defined tortuosity as an empirical coefficient to describe the random orientation of pores. Porosity and tortuosity in porous glass membranes were described using percolation theory. Tortuosity factors were determined not only by diffusion paths but also by the amount of throughout porosity. If the internal geometry of the porous glass membrane is known, tortuosity can be calculated by Monte Carlo simulations.9 Tortuosity factors were calculated on the basis of the porosity (fraction of leachables). For low porosities (20). This was explained by means of the percolation threshold. For porous VYCOR glass (porosity 31%), a tortuosity of 6.5 was calculated from theory in good accordance with the experimental value of 5.9. The theoretical calculations also showed that an increase in the porosity of the membranes up to 60% results in a decrease of the tortuosity to a value of 1.5. 60 nm CPG membranes as those used here are characterized by a porosity of approximately 50%. Additionally, tortuosities of porous glass membranes of the same type were obtained from measurements of the permeability of nitrogen.10 Several porous glass membranes with 0.24 mm thickness, pore sizes between 2 and 20 nm, and porosities between 20 and 60% were examined for the permeability of nitrogen at 298 K and 1.1−1.5 bar. The measurements were performed as follows: the membrane was stuck on a brass plate containing a bore with 2 and 5 mm diameter, respectively. The plate was fixed with a special dome on a turbo molecular pump (TMU 261, Pfeiffer). The dome was evacuated. The pumping speed was 210 L s−1 (as specified by the manufacturer). The gas flow was determined by measuring the pressure below the membrane. Finally, the integral permeability was estimated using the pumping speed, the measured pressures above and below the membrane, the cross-section of the bore, and the membrane thickness. Here, the pumping speed was independent of the pressure under the experimental conditions. The tortuosity factors were calculated on the basis of (i) integral permeability, (ii) porosity, and (iii) the Knudsen diffusion coefficient of nitrogen at 298 K in a straight cylindrical pore with the corresponding pore diameter.10 In this study tortuosity factors of 1.5 and 1.6 were observed for membranes with 10, 12, and 18 nm pore diameter and porosities between 50 and 60%. Sample Preparation. Prior to infiltration, AAO and CPG membranes were annealed at 180 °C for 2 h under a vacuum. CPG was infiltrated with ACE in a beaker placed on a hot-stage at 180 °C for 2 min under ambient conditions. Subsequently, the infiltrated CPG monoliths were placed in furnaces, heated to 175 °C, and subjected to non-isothermal crystallization at a cooling rate of 0.5 K/min under argon in the presence of bulk 80



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scattering intensity profiles I(Ψ) along Debye rings (Figure S1, on the right) belonging to the fixed scattering angles Θ. The Debye ring is the circle of intersection of the Ewald sphere and a sphere about the origin of the reciprocal space, the radius of which is the length of the scattering vector belonging to Θ. Thus, I(Ψ) represents the orientation distribution of a set of lattice planes belonging to the selected scattering angle Θ relative to the AAO surface. Likewise, I(Ψ) represents the orientation distribution of the corresponding reciprocal lattice vectors with respect to the AAO pore axes. Note that the apparent I(Ψ) values sharply decrease for high Ψ angles > ∼70° owing to defocusing effects.12 Hermans’ order parameter f13 can be calculated from the I(Ψ) profiles according to

f=

3 1 ⟨cos2 ψ ⟩ − 2 2

Figure 3. First DSC heating scans of ACE non-isothermally crystallized at a rate of −0.5 K/min in the presence of bulk ACE surface films in AAO and CPG, followed by removal of the bulk ACE surface films and subsequent annealing at 135 °C for 2 h (AAO) or 1 h (CPG).

(1)

where k

⟨cos2 ψ ⟩ =

∑i = 1 [I(ψ )i cos2 ψi ] k

∑i = 1 I(ψ )i

We investigated the AAO sample of which a DSC heating run is displayed in Figure 3 by wide-angle X-ray scattering. Using Θ/2Θ geometry allowed detection of reflections originating from sets of lattice planes oriented normal to the AAO pore axes (parallel to the AAO surface, see Methods and Figure S1). Figure 4a shows a Θ/2Θ scan taken still in the

and k is the number of data points of the Schulz scan, that is, the number of angles Ψi at which I(Ψ) was measured.

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RESULTS AND DISCUSSION Results. Three polymorphs of ACE have been identified. Monoclinic form I of ACE has a bulk melting temperature T m∞,I ∼ 167−169 °C14,15 and can be obtained inside CPG by quenching from the melt while in contact with a bulk ACE surface reservoir. Form III16,17 (bulk melting temperature Tm∞,III ∼ 143 °C) is inaccessible in bulk systems but forms in CPG with pores smaller than ∼100 nm by isothermal crystallization in the absence of bulk ACE surface reservoirs.18 Orthorhombic form II15 with a bulk melting temperature Tm∞,II = 156−158 °C14 cannot be produced inside CPG with pores smaller than 100 nm except by subjecting form III crystals to tedious thermal cycling procedures.19 When confined to nanoporous hosts, all three polymorphs show pronounced melting point depression related to the Gibbs−Thomson effect, which can be estimated using a modified Gibbs−Thomson equation.20,21 Figure 3 in ref 18 shows Gibbs−Thomson plots (melting temperature as a function of the inverse pore diameter) of forms I, II, and III. Strikingly, non-isothermal crystallization in contact with bulk ACE surface films at a cooling rate of −0.5 K/min resulted in predominant formation of form I in CPGs and of form II in AAO. Occasionally, the AAO samples also contained form III that could be converted to form II by annealing at 135 °C. This temperature is above the melting point of form III and below the melting point of form II of ACE confined to 60 nm pores.18 DSC heating scans of ACE in AAO (Figure 3) confirmed the presence of form II, as obvious from the position of the melting peak (Tm,onset = 150 °C, Tm,peak = 153 °C), which is in line with the position expected for form II in 60 nm nanopores.18 The presence of a minor fraction of form III is indicated by a melting peak at Tm,peak = 137 °C. In CPGs, however, the melting peak (Tm,peak = 165 °C) is shifted to temperatures where form I is expected to melt when confined to 60 nm pores. A small melting endotherm at Tm,peak ∼ 169 °C can be ascribed to residual bulk ACE form I on the CPG surface (the bulk ACE surface films had been scraped off before any measurement).

Figure 4. X-ray analysis of ACE (dark yellow) in AAO (black) cooled at a rate of −0.5 K/min in the presence of a bulk ACE surface film (red). All patterns were taken from the sample of which the first DSC heating scan is shown in Figure 3 and all main reflections could be ascribed to form II. (a, b) Θ/2Θ scan taken in the presence of a bulk ACE surface film directly after non-isothermal crystallization. In panels (a) and (b), the intensity axis is linear but rescaled in (b) so that the weak peaks are more discernible while the strong peaks are cut off. The X-ray patterns of (a) and (b) are dominated by crystals located in the bulk ACE surface film. (c) Θ/2Θ scan taken after removal of the bulk ACE surface film and annealing at 135 °C for 2 h. The Θ/2Θ scan taken after removal of the bulk ACE surface film before the additional annealing step at 135 °C was identical. The detected scattering intensity exclusively originated from ACE located inside the AAO pores. (d) Schulz scans taken after removal of the bulk ACE surface film and annealing at 135 °C for 2 h.

presence of the bulk ACE surface film (red top layer in Figures 1b and 4) directly after non-isothermal crystallization so that the diffraction pattern was dominated by the bulk ACE surface film. Several intense higher-order (00l) reflections of form II appeared, which indicated the presence of ACE form II crystals having their (00l) lattice planes oriented parallel to the AAO surface in the bulk ACE surface film (form II reflections were indexed according to Cambridge Structural Database 22 deposition number CSD-HXACAN23 based on ref 23). 81



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Rescaling the intensity axis of Figure 4a reveals the presence of weak form II (020) and (040) reflections as well as of a weak form II (120) peak (Figure 4b). No reflections that could not be ascribed to form II appeared. Next, we measured Θ/2Θ scans after removal of the bulk ACE surface film. Hence, the detected scattering intensity exclusively originated from ACE located inside the AAO pores (dark yellow in Figures 1b and 4). The X-ray patterns thus obtained showed intense (020) and (040) reflections of form II at 2Θ ≈ 14.9° and 2Θ ≈ 30.1° (Θ/ 2Θ patterns measured directly after removal of the bulk ACE surface film and after an additional annealing step at 135 °C for 2 h without bulk ACE surface film were identical). All other reflections were suppressed (Figure 4c). This outcome indicated uniform orientation of the ACE form II crystals located in the AAO pores on a macroscopic scale; their (0k0) faces were oriented parallel to the AAO surface and normal to the AAO pore axes. We could unambiguously identify form II because powder patterns of form I do not show significant reflections around 2Θ ≈ 30°, and powder patterns of form III do not show significant reflections around 2Θ ≈ 15°.16,19 The apparent crystal texture was further investigated by Schulz scans11,24 (cf. Methods and Figure S1) taken after removal of the bulk ACE surface film and annealing for 2 h at 135 °C (Figure 4d). Schulz scans yield azimuthal intensity distributions I(Ψ) (Ψ is the azimuthal angle) along the Debye rings belonging to fixed scattering angles Θ. The I(Ψ) profiles represent, therefore, orientation distributions of sets of lattice planes belonging to specific reflections relative to the AAO surface (or orientation distributions of the corresponding reciprocal lattice vectors with respect to the AAO pore axes). The Schulz scans confirmed the strong alignment of the form II (0k0) faces with the AAO surface (or the strong alignment of the reciprocal (0k0) lattice vectors with the AAO pore axes), as apparent from narrow I(Ψ) maxima at Ψ = 0° for the (020) and (040) reflections. The Hermans order parameters (Methods, eq 1)13 amounted to ∼0.97. The occurrence of the I(Ψ) maximum at Ψ ∼ 39° for the (120) reflection of ACE form II (2Θ ∼ 19.1°) (Figure 4d) is consistent with strong alignment of the (0k0) lattice planes with the AAO surface as the angle between the (0k0) and (120) lattice planes is 38.6°. Directly after the non-isothermal crystallization step, a vast majority of AAO samples exclusively showed intense (020) and (040) reflections of form II. Occasionally, after the nonisothermal crystallization step form III dominated inside the AAO pores. Figure 5 shows Θ/2Θ scans of such a sample directly after non-isothermal crystallization in the presence (Figure 5a) and after removal (Figure 5b) of the bulk ACE surface film as well as after an additional annealing step at 135 °C for 2 h (Figure 5c). In Figure 5a several form III reflections predominantly originating from crystals in the bulk ACE surface film can be seen. After removal of the bulk ACE surface film, the Θ/2Θ pattern is dominated by a strong form III (400) reflection at 2Θ ≈ 30.1°. This outcome suggests that the AAO pores contained uniaxially oriented form III ACE crystals having their (400) lattice planes oriented normal to the AAO pore axes. Annealing at 135° for 2 h converted the uniaxially oriented form III crystals in the AAO pores into likewise uniaxially oriented form II crystals having their (0k0) lattice planes oriented normal to the AAO pore axes. To study isothermal crystallization, ACE-infiltrated AAO covered by a bulk ACE surface film was quenched from 180 °C to below −10 °C. Subsequent isothermal crystallization in the presence of a bulk ACE surface film at 80 °C for 2 h yielded

Figure 5. X-ray analysis of ACE (dark yellow) in AAO (black) cooled at a rate of −0.5 K/min in the presence of a bulk ACE surface film (red). (a) Θ/2Θ scan taken in the presence of a bulk ACE surface film directly after nonisothermal crystallization. (b) Θ/2Θ scan taken after non-isothermal crystallization and subsequent removal of the bulk ACE surface film. (c) Θ/2Θ scan taken after non-isothermal crystallization, removal of the bulk ACE surface film, and annealing at 135 °C for 2 h.

form III. A Θ/2Θ scan taken after removal of the bulk ACE surface film showed a dominating form III (400) reflection 2Θ ≈ 30.1°. The appearance of a weak form II (200) peak at 2Θ ≈ 14.9° indicated the presence of a minor form II fraction (Figure 6a). The absence of any other reflections indicated pronounced crystal orientation as in the case of the non-isothermally crystallized samples; form III (h00) faces and form II (0k0) faces were preferentially oriented normal to the AAO pore axes. Annealing at 135 °C for 2 h after removal of the bulk ACE surface film resulted in the conversion of form III into form II. Θ/2Θ patterns measured after annealing at 135 °C (Figure 6b) showed inverted relative intensities of the peaks at 2Θ ≈ 14.9° and 2Θ ≈ 30.1°; the Θ/2Θ pattern showed a dominant form II (020) peak at 2Θ ≈ 14.9° and corresponded to the form II patterns seen in Figures 4c and 5c. Orientation distributions revealed by Schulz scans at 2Θ ≈ 30.1° corresponding to the form III (400) peak and the form II (040) peak by and large coincided before and after the conversion of form III to form II (Figure 6c). The pronounced uniaxial orientation of both form III and form II crystals in the AAO pores, as obvious from narrow I(Ψ) maxima at Ψ = 0° and Hermans order parameters larger than 0.95, persisted after the solid/solid transition from form III to form II. Schulz scans taken prior to the solid/solid transition for a 2Θ value of ∼19.1° represent orientation distributions of form III (211)/(202) faces with respect to the AAO surface, whereas Schulz scans for the same 2Θ value taken after the solid/solid transition represent orientation distributions of form II (120) faces with respect to the AAO surface. The angles between form III (211) and (202) faces on the one hand and form III (h00) faces on the other hand are ∼39°, such as the angle between form II (120) faces and form II (0k0) faces. The I(Ψ) profiles obtained for 2Θ = 19.1° before and after the solid/solid transition showed coinciding peaks at Ψ ≈ 39°. Again, the solid/solid transition of form III into form II did not result in a significant change of the peak shape. As discussed above, exclusively form I melting peaks appeared in DSC traces of ACE located in CPG that was crystallized at a cooling rate of −0.5 K/min in the presence of a bulk ACE surface film (cf. Figure 3). Additional heating to 135 °C for 1 h without bulk ACE surface film did not change the 82



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Figure 6. Isothermally crystallized ACE in AAO with a pore diameter of 60 nm. (a) Θ/2Θ scan taken after cold crystallization at 80 °C for 2 h in the presence of a bulk ACE surface film and removal of the bulk ACE surface film; (b) Θ/2Θ scan taken after subsequent annealing at 135 °C for 2 h. (c, d) Schulz scans taken for 2Θ angles of (c) 30.1° and (d) 19.1°. Red Schulz scans correspond to the Θ/2Θ scan of (a) and were taken directly after crystallization at 80 °C and removal of the bulk ACE surface film. Black Schulz scans correspond to the Θ/2Θ scan of (b) and were taken after additional heating to 135 °C for 2 h after removal of the bulk ACE surface film.

AK6 yielded calculated powder patterns that closely matched observed powder patterns of form III. However, since form III readily undergoes an apparently monotropic solid/solid transition to form II,17,27 Coombes et al. suggested that the structures of both polymorphs may be more closely related.28 Moreover, several peaks coincide in the powder patterns of form II and form III. This finding was attributed to matching unit cell parameters of both forms consisting of topologically equivalent two-dimensional layers of hydrogen-bonded ACE molecules.17 The extension of a proposed orthorhombic form III unit cell in the direction equals the extension of the likewise orthorhombic form II unit cell in the direction (Figure 3 in ref 17). Non-isothermal crystallization of ACE in AAO with bulk ACE surface film at a cooling rate of −0.5 K/min occasionally yielded form III with uniform crystal orientation. Likewise, the cold crystallization experiment summarized in Figure 6 (crystallization of vitrified ACE in AAO at 80 °C for 2 h in the presence of a bulk ACE surface film) yielded uniformly oriented form III crystals. Annealing at 135 °C readily transformed the form III crystals in the AAO pores having their (h00) faces oriented normal to the AAO pore axes to form II crystals having their (0k0) faces oriented normal to the AAO pore axes. This result indeed suggests a high degree of structural registry between the form III direction and the form II direction (which are oriented normal to the respective lattice planes as the unit cells are orthorhombic). In CPG form III can be produced by cold crystallization in the absence of bulk ACE surface films (in the presence of bulk ACE surface films form I forms even at high cooling rates). However, inside CPG the transition from form III to form II is suppressed; high yields of form II are only accessible by tedious thermal cycling procedures.19 The striking dependence of the kinetics of the solid/solid transition from form III to form II on the pore morphology might have two reasons. First, the transition front might rapidly propagate along the form III direction because of structural matching of form III along the form III direction and form II along the form II direction. The propagation paths along straight AAO pores are 3 orders of magnitude longer than those within the spongelike pore systems of CPG (∼100 μm vs ∼100 nm). Whereas the front of the transition from form III to form II can propagate through AAO pores without being impeded by obstacles, in CPG the propagation front will impinge on pore walls after propagation paths of ∼100 nm (in analogy to the

melting behavior of the ACE (cf. Figure 3). Typically, corresponding Θ/2Θ scans (Figure 7a) showed reflections

Figure 7. ACE form I in CPG and AAO. (a) Exemplary Θ/2Θ scan of ACE non-isothermally crystallized in CPG at a rate of −0.5 K/min in the presence of a bulk ACE surface film measured after removal of the latter and annealing at 135 °C for 1 h; (b) Θ/2Θ scan of ACE in AAO obtained by quenching from 175 °C to below 0 °C in the presence of a bulk ACE surface film measured after removal of the latter. Selected form I reflections are indicated.

compatible with form I (indexed according Cambridge Structural Database deposition number CSD-HXACAN03 based on ref 25) with relative intensities deviating from those apparent in powder patterns. Inside the continuous CPG pore networks larger crystals dominated, as obvious from the fact that in corresponding X-ray patterns typically only a few reflections appeared. However, the X-ray patterns of different samples showed different relative peak intensity profiles, suggesting that the large form I crystals in the volume illuminated by the X-ray beam had random orientations with respect to the scattering geometry. Quenching ACE in AAO in the presence of a bulk ACE surface film from 175 °C to below 0 °C by placing the sample on a cooled copper plate yielded form I, as confirmed by DSC measurements (Figure S2a, Supporting Information). The Θ/ 2Θ scan measured after removal of the ACE surface film (Figure 7b) contained form I reflections with relative intensities deviating from those of powder patterns. Schulz scans (Figure S2b) revealed the coexistence of several populations of ACE form I crystals having different orientations. Influence of Pore Morphology on the Form III/Form II Solid/Solid Transition. Peterson et al.16 reported that a monoclinic structure predicted by Beyer et al.26 referred to as 83



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orientation requires a certain residence time at low supercooling. A specific nucleation process effective at low supercooling similar to that previously proposed for the crystallization of polyethylene in AAO47 appears to be central to the preferred growth of specifically oriented form II or form III crystals in AAO. Bulk ACE crystals formed by heterogeneous nucleation in bulk ACE surface films may impinge on the AAO surface and offer pre-existing crystal surfaces for sporadic heterogeneous nucleation at the mouths of the AAO pores (Figure 1a). Heterogeneous nucleation of ACE was reported to be most efficient if surface functionalities rather than lattice parameters of the counterpart crystal faces match.48 Alternative nucleation mechanisms that initiate the formation of ACE crystals in the AAO pores away from the pore mouths must be suppressed as ACE crystals originating from these alternative nucleation mechanisms would clog the growth paths of crystals formed close to the pore mouths. As discussed above, we can rule out surface nucleation at the AAO pore walls as well as heterogeneous nucleation processes inside the AAO pores other than heterogeneous nucleation at the pore mouths. (2). Competition Between Different Growing Crystal Species. During crystal growth growing crystals compete for growth volume. It is reasonable to assume that pronounced differences in the growth rates along directions normal to specific crystal faces cause the dominance of specifically oriented crystals in AAO pores. Crystal growth rates along certain directions are related to attachment rates of molecular building blocks to crystal faces normal to the growth directions. The Hartman−Perdok model distinguishes between rough growth on rough crystal faces and layered growth on flat crystal faces.49 Flat crystal faces are characterized by the presence of connected nets of the molecular building blocks of the crystal, whereas such connected nets of reasonable strength are absent on rough crystal faces.50 Rough crystal growth normal to rough crystal faces is orders of magnitude faster than growth normal to flat crystal faces.51 Heterogeneous nucleation at bulk ACE crystals impinging on the AAO surface may yield ACE crystals of several polymorphs with several orientations. Slowly growing crystals or crystals having their fast growth direction not aligned with the AAO pore axes (so that they impinge on the AAO pore walls) do not significantly contribute to the crystallization of the ACE located in the AAO pores and may even serve as substrates for further heterogeneous nucleation events. However, crystals that have rough crystal faces without connected nets of molecular crystal building blocks oriented normal to the AAO pore axes have their direction of fast growth aligned with the AAO pore axes; these crystals will rapidly grow along the AAO pores until they occupy the entire pore volume. The cylindrical hard confinement imposed by the rigid AAO pore walls suppresses the transformation of rough crystal faces into edges or vertices, as commonly observed for rough crystal faces in the course of bulk crystal growth,52 and stabilizes steady-state growth normal to the rough crystal faces. (3). Influence of Free Path Length. Ultimately, the preferential formation of different ACE polymorphs in AAO and CPG can be related to the different morphologies of the host systems. It is reasonable to assume that form III (h00) faces exist that do not exhibit connected nets and that show rough growth within the range of crystallization temperatures considered here. Once formed, form III crystals having their direction aligned to the AAO pore axes rapidly grow down the AAO pores. The subsequent solid/solid transition

situations illustrated in Figure 1b,c). Second, the solid/solid transition from form III to form II in the continuous isotropic pore systems of CPG would lead to the formation of form II grains having differing orientations that would presumably be separated by energetically unfavorable grain boundaries. Influence of Pore Morphology on Crystallization. Uniaxial crystal growth was studied theoretically29−32 and experimentally. Various mostly polymorphic organic and polymeric materials show pronounced oriented crystallization in straight cylindrical nanopores. Examples include semicrystalline polymers,33−35 n-hexane,36 longer-chain n-alkanes,37 linear 1-alcohols,38 ROY,39 and glycine.40 In some cases different polymorphs were obtained by different thermal treatments,34 and occasionally hosts with different pore morphologies were used.41 For example, in mesoporous silicon with straight aligned mesopores n-alkanes formed lamella crystals. The molecular long axes of the n-alkanes in the lamella crystals were oriented normal to the mesopore axes.37 However, in Vycor glass having randomly oriented, spongelike pore systems (similar to CPG), lamella formation was suppressed, and the n-alkanes formed a nematocrystalline phase.42 As far as we are aware, systematic evaluations of the influence of the morphologies of porous hosts on the crystallization of polymorphic materials have not been reported. On the basis of the results summarized above, we propose a qualitative model for the morphology dependence of the crystallization of polymorphic materials in nanoporous hosts. Three aspects play a key role: (1) the way in which crystallization is initiated by nucleation; (2) the competition between different crystal species for growth volume; and (3) the length of the free growth path available to growing crystals. (1). Influence of the Nucleation Mechanism. Nucleation initiates crystallization. Heterogeneous nucleation43,44 is catalyzed by solid heterogeneities such as impurities, dust, nucleating agents, or other crystals that expose solid surfaces to supercooled or vitrified liquids. Germs of the crystallizing species more easily form at surfaces of solid heterogeneities than in metastable liquids away from solid heterogeneities. The contacts to solid heterogeneities reduce the germs’ overall surface energies and in turn their critical radii. Therefore, heterogeneous nucleation sets in at significantly lower supercooling than alternative nucleation mechanisms. Homogeneous nucleation,45 the formation of nuclei without involvement of heterogeneities, occurs at lower temperatures than heterogeneous nucleation; during cooling, homogeneous nucleation is typically characterized by formation of many homogeneous nuclei within a narrow temperature window.46 Heterogeneous nucleation can be suppressed by confining the crystallizing material to small containers such as AAO pores unlikely to contain heterogeneous nuclei and by rapid cooling from the melt so that homogeneous nucleation sets in before growing crystals originating from heterogeneous nucleation have consumed significant portions of the metastable melt. In the case of ACE non-isothermally crystallized in AAO in the presence of a bulk ACE surface film at −0.5 K/min, a mechanism must exist that yields uniformly oriented form II or form III crystals having their form II (0k0) or form III (h00) faces oriented perpendicularly to the AAO pore axes, even though the crystals in the bulk ACE surface films may have different orientations. Moreover, rapid quenching suppresses formation of uniformly oriented form II/form III crystals in the AAO pores and yields form I crystals having several different orientations. Apparently, kinetic selection of a specific growth 84



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polymorphs than in straight cylindrical pores. The solid/solid transition from form III to form II readily occurs in the straight cylindrical AAO pores because the transition front can propagate through the AAO pores, along which crystals of both involved polymorphs having the preferentially occurring orientation show high structural registry. In CPG, the propagation front again impinges on CPG pore walls after short propagation paths. Moreover, in CPG formation of form II crystals having different orientations might be accompanied by formation of energetically disadvantageous grain boundaries. Pore morphologies of nanoporous hosts are an important parameter for the control of confined crystallization and of solid/solid transitions of organic and polymeric materials in nanoporous host systems. Moreover, the results presented here should have a significant impact on mesoscopic crystal engineering of drug delivery systems for rational drug release management. For example, polymorphs otherwise not accessible in nanopores can be produced at high yields. Investigating crystallization of drugs in AAO is of particular interest as replication of AAO has been explored as an attractive route to nanoscale cylindrical drug carriers.53,54 Strategies for mesoscopic crystal engineering in AAO pores can be transferred to other potential nanoscopic drug carrier systems. For example, AAO can be regarded as a model system for arrays of titania nanotubes, which are accessible by atomic layer deposition into block copolymer templates and subsequent removal of the block copolymer.55 It should be straightforward to transfer the mesoscopic crystal engineering principles elaborated here to other areas such as nanowire-based organic electronics.56

propagating along the AAO pores without being impeded by obstacles yields highly oriented form II crystals having their direction aligned with the AAO pore axes. Despite the pronounced structural registry of correspondingly oriented form III and form II crystals along the AAO pore axes, form III (h00) faces and form II (0k0) faces oriented normal to the AAO pore axes are different in their growth characteristics. The assumption that form III is the actual growth form while form II is obtained from form III via an apparently monotropic solid/ solid transition leads to the conclusion that no form II face without connected nets showing rough growth exists that can kinetically compete with form III (h00) faces. Presumably, the same crystallization mechanism than in AAO takes place in CPG too, yielding ACE crystals having their directions of fast crystal growth oriented normal to the pore axes. However, the CPG pores are curved; the free growth paths in the meandering CPG pores are 3 orders of magnitude shorter than those in AAO pores (∼100 nm vs ∼100 μm). As schematically displayed in Figure 1c, in CPG fast-growing ACE crystals inevitably impinge on CPG pore walls. The limited free growth path in CPG for any fast-growing crystal species prevents the kinetic selection process effective in AAO and allows domination of other modes of crystallization under otherwise identical crystallization conditions. The persistence of form III obtained in CPG by isothermal cold crystallization at 80 °C18 appears to be a kinetic phenomenon.

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CONCLUSIONS Using acetaminophen as an example, we have shown that under otherwise identical conditions (i) molecular crystals can preferentially form different polymorphs in nanoporous hosts with different pore morphologies and that (ii) solid/solid transitions between different polymorphs are affected by pore morphology. In the presence of a bulk ACE surface reservoir, non-isothermal crystallization of ACE in the straight cylindrical pores of self-ordered AAO (tortuosity = 1) at slow cooling rates yields uniformly oriented form II and/or form III crystals. The occurring orientations of form II and form III crystals are characterized by high structural registry along the AAO pores. Under otherwise identical conditions, form I crystals preferentially form in CPG with continuous, isotropic pore systems (tortuosity >1.5). The uniformly oriented form III crystals in AAO readily undergo a solid/solid transition to likewise uniformly oriented form II crystals that can be produced at high yield in AAO. The solid/solid transition from form III to form II is suppressed in CPG. We suggest that these findings can be rationalized as follows. Bulk crystals located in surface films on the AAO membranes impinging on the AAO provide surfaces for sporadic heterogeneous nucleation at the pore mouths. In this way, occasionally crystals having a distinct direction of fast crystal growth aligned with the AAO pore axes may form and grow rapidly along the AAO pores (free growth path in AAO ∼ 100 μm). Thus, polymorphs having a distinct direction of fast crystal growth will dominate in AAO pores, and the crystals of this polymorph will preferentially be oriented in such a way that crystal faces with high attachment rates are oriented normal to the AAO pore axes. Free growth paths in tortuous pore systems, such as the pore systems in CPG, are orders of magnitude shorter than in the straight cylindrical pores of AAO. Therefore, fast-growing crystals impinge on the CPG pore walls and occupy only small volume fractions, while alternative modes of crystallization prevail in CPG and possibly result in dominance of other

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ASSOCIATED CONTENT

S Supporting Information *

Figure S1: Scheme of the setup used for X-ray analysis; Figure S2: DSC scan and texture analysis of ACE form I in AAO obtained by rapid cooling. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank the German Research Foundation (BE 2352/4, EN 942/1, STE 1127/14, INST 190/134-1 FUGG) for financial support, C. Hess and H. Tobergte for the preparation of AAO, and S. Sklarek (MPI of Microstructure Physics, Halle) for additional SEM investigations. We wish to acknowledge the use of the EPSRC funded Chemical Database Service at Daresbury.

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REFERENCES

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Morphology of Porous Hosts Directs Preferred Polymorph Formation

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