Rheology of Cellulose Ether Excipients Designed for Hot Melt Extrusion

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Rheology of Cellulose Ether Excipients Designed for Hot Melt Extrusion Tirtha Chatterjee, Kevin P O'Donnell, Mark A Rickard, Brian Nickless, Yongfu Li, Valeriy V. Ginzburg, and Robert L. Sammler Biomacromolecules, Just Accepted Manuscript • DOI: 10.1021/acs.biomac.8b01306 • Publication Date (Web): 15 Oct 2018 Downloaded from http://pubs.acs.org on October 22, 2018

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Rheology of Cellulose Ether Excipients Designed for Hot Melt Extrusion

Tirtha Chatterjee*1, Kevin P. O’Donnell2, Mark A. Rickard3, Brian Nickless4, Yongfu Li3, Valeriy V. Ginzburg4, Robert L. Sammler4

1

Dow Water & Process Solutions, The Dow Chemical Company, Collegeville, PA 19426 2 3

4

Dow Food, Pharma & Medical, The Dow Chemical Company, Midland, MI 48674

Analytical Sciences, Core R&D, The Dow Chemical Company, Midland, MI 48667

Materials Science and Engineering, Core R&D, The Dow Chemical Company, Midland, MI 48674

1 ACS Paragon Plus Environment

Biomacromolecules 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Abstract

A new family of cellulosic ether polymeric excipients has been recently engineered for fabrication of amorphous solid dispersions (ASDs) of active pharmaceutical ingredients (API) via hot-melt extrusion (HME). These hydroxypropyl methyl cellulose (HPMC) excipients enable plasticizer-free melt processing at much lower temperatures (135 to 160 °C) due to their substantially-reduced glass transition temperatures Tg (98 to 110 °C). The novel amorphous cellulose ethers were found to be rheologically solid-like well above their glass transition (Tg + 70 °C). We demonstrate that in pharmaceutically-relevant HME processing temperature range these polymers behave similarly to yield-stress fluids and flow only when the applied stress exceeds a critical stress value. This critical stress value (0.50 ± 0.05 MPa, 150 °C) is surprisingly high but is easily achieved at typical HME conditions. The origin of their yield stress fluid-like behavior is hypothesized to arise from hydrogen bonds of the HPMC materials that act as physical crosslinks and do not relax within the measured temperature and time window unless the applied stress exceeds the critical stress value. Supports for this hypothesis arises from infrared spectroscopic estimates of the free- and bound- hydrogen bond levels at end-use temperatures.


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Introduction Active pharmaceutical ingredients (APIs) are classified based on their permeability and solubility according to the biopharmaceutics classification systems (BCS). A large section of the marketed APIs fall under either class II (high permeability and low solubility) or class IV (low permeability and low solubility) categories. Furthermore, the majority of the new chemical entities in development as potential APIs are poorly water soluble.1 Formulation into an amorphous solid dispersion (ASD) is one of the leading strategies to enhance an API’s dissolution rate and extent through dispersing them in their amorphous form in an excipient matrix.2 The apparent solubility of a true amorphous API is orders of magnitude higher than the thermodynamic solubility of its crystalline counterpart. The two most common industrially-practiced ASD processing technologies are spray drying (spray-dried dispersion, SDD)3-4 and hot melt extrusion (HME).5-6 In SDD a feed solution of the polymer and API is atomized and rapidly dried to maintain the active in an amorphous state. The SDD technology is widely practiced industrially in spite of some limitations such as: (a) selected choice of organic solvents available for API and polymer co-solubilization; (b) solvent cost, flammability, and toxicity; (c) the need of post-processing to remove residual solvents in the end-product below the regulatory limit; (d) significant physical footprint and finally, (e) challenges associated with spraying a viscous feed when prepared with high molecular weight polymeric excipients. Recently, hot melt extrusion has grown significantly in the pharmaceutical industry as a solvent-free alternative approach to ASD manufacturing. Amorphous solid dispersions are prepared by warming and mixing the molten excipient and API (inside an extruder barrel) followed by rapid cooling to a vitrified state at the extruder die exit. The timescale of vitrification/solidification must be orders of magnitude faster than the API recrystallization timescale to ensure their dispersion in amorphous state. This is a robust, scalable, continuous processing technique which is approved by the US Food and Drug Administration (FDA). The advantages and limitations of the HME-based ASD has been covered in several recent publications.5-15



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Hydroxypropyl methyl cellulose (HPMC) has been a key polymeric excipient for delivery of actives in both immediate- and controlled-release applications.16 Henceforth, the term polymeric excipients will be referred to as polymer for brevity. These polymers were not originally suitable for HME processing without a significant level of plasticizer addition for most actives17-18 as their glass transition is close to or above their caramelization temperature (typically ≈ 170 °C). The HME polymers must be thermally stable and processable between the temperature at which the API can enter the matrix (this can be above or below the melting temperature of the API) and the API degradation (Td) temperature.6 Recently, a special class of HPMC polymers19 have been developed for actives with poor solubility. Specifically, a novel grade10, 20 has been commercialized that possess a glass transition temperature (Tg) suitable for pharmaceutical HME applications without using any plasticizer.17, 20 Hereafter, this novel grade will be referred to as HPMC HME polymers. The low Tg is hypothesized to result from the high level of substitution of methyl ether (MeO) and hydroxypropyl ether (HPO) side chains grafted to the anhydroglucose backbone. This high substitution level is also beneficial in impeding API recrystallization under storage conditions, and increasing API’s water solubility and potentially bio-availability when delivered in the gastrointestinal tract.10 The two most critical physical properties relevant for HME are carrier polymer thermal stability and rheology under relevant processing conditions. Commercial HPMC polymers, including HPMC HME grades, are relatively well studied in terms of thermal stability18, 21-22 while limited information20, 23 is available on their rheology. The carrier polymer rheology enables the formulators and process engineers to select and optimize extrusion conditions (temperature, shear rate range, screw design etc.). The polymer rheology data are also required to model blend rheology profiles once an API is incorporated. Motivated by these needs, the current study is focused on novel HPMC HME polymer melt rheology. We demonstrate that in pharmaceutical-relevant HME processing conditions these polymers exhibit yield-stress fluid-like rheology and flow only when the applied stress exceeds a critical, temperature dependent, stress value. The critical stress values under uniaxial tensile deformation has been measured for the HPMC polymers suitable


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for HME processing. This critical stress value was found to be surprisingly high, inaccessible to commercial rotational rheometers (for true melt-state viscoelastic properties measurement) but could easily be achieved at typical extrusion conditions to prepare a homogeneous molten API/polymer dispersion. Finally, a possible origin behind this unique rheological behavior has been advanced.

Materials and methods Materials: Three different grades of commercial HPMC HME materials used in this study are the commercial products of The Dow Chemical Company which are available under the trade name AFFINISOL™.i These polymers are commercially denoted as AFFINISOLTM HPMC HME 15-, 100-, and 4M cP polymers and differ by their molecular weight. More details of these polymers are reported in Table 1. Sample names are after their 2% solution viscosity in water at 5 C as typically practiced in pharmaceutical industry. It is evident from Table 1 that the polymer chain molecular weight (weightaverage, Mw) is inversely related to their solution viscosity value. Also these three grades have similar substitution levels within the range reported in Table 1.

Sample Preparation: The HPMC powders were dried overnight under vacuum at 80 °C. The compression molding was performed by pre-heating the dried powders to 140 °C and holding at that temperature for 2 min under a compressive force of 30,000 kg followed by cooling down to room temperature (without removing the applied force) with a typical cool down time of ~10 min. Circular or rectangular shaped specimens of desired dimensions were cut from the compression-molded plaques for rheological testing. Dog bone shaped specimens (tensile bars) were used for the uniaxial deformation stress-strain measurements. A total of 5 tensile bar specimens were molded during each cycle. Specimens were stored in a desiccator prior to use to minimize water absorption and avoid bubble formation. Molecular

i

™ Trademark of The Dow Chemical Company (“Dow”) or an affiliated company of Dow



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weights of the polymers collected from the plaques were measured and compared with the powder (Table 1) to assess whether the molding process altered the long-chain structure of the polymers.

Rheology: A strain-controlled ARES-G2 melt rheometer (TA Instruments, TRIOS data acquisition software, version 4.3.1.39215) was used for the majority of the rotational rheology measurements reported here. All ARES-G2 measurements were performed using a 25-mm parallel plate fixture (material of construction: stainless steel). A strain-controlled ARES rheometer (TA Instruments, ORCHESTRATOR data acquisition software, version 6.04.00) was used for temperature sweep measurements using a torsional bar fixture. All experiments were conducted under N2 environment to minimize/avoid thermal oxidation of polymer at elevated temperatures. A fresh aliquot was used to measure a series of test sequences: the strain amplitude sweep followed by dynamic frequency sweep both under oscillatory shear. The applied oscillatory shear strain range in the amplitude sweep was 0.01-100%. Both the real (G’) and imaginary (G”) components of the complex shear modulus (G*) were monitored as a function of strain amplitude at a fixed angular oscillation frequency (ω) of 1 rad/s to determine the linear viscoelastic (LVE) regime. Components G′ and G″ are also known as the shear storage and shear loss modulus, respectively. By definition, in the LVE regime, the moduli are independent of the applied strain amplitude. Dynamic frequency sweep (FS) was subsequently performed for an angular frequency ω range of 398.1 to 0.1 rad/s with the applied oscillatory strain amplitude in the LVE regime. Both the G’ and G” were monitored as a function of frequency. The magnitude component of the complex shear viscosity was derived as: |η*|=|G*|/ω, where the amplitude component of the complex shear modulus is given as |G*| = (G’2+ G” 2)1/2. The typical relative precision (2 std. deviation/mean, in %) of the melt viscosity measured by the rotational rheometer is < 5%. For temperature sweep measurements, a rectangular cross-sectional specimen was mounted in a torsional bar fixture and the sample was heated from -80 to 200 °C at a warming rate of 10 °C /min. An oscillatory strain with amplitude and angular frequency 0.25% and 1 rad/s, respectively was used. Both the G’ and G” were monitored as a function of temperature. The tangent of the phase angle () is defined as tan = G”/G’.


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The tan vs. temperature data were fitted to a Gaussian model on linear axes to obtain the peak temperature assigned as the glass transition temperature. An Instron frame model 5567A with a static load cell of 100N was used to generate the engineering tensile stress-strain data. The test specimen had an overall length of 3” (7.62 cm). The width of the narrow section was 0.125” (~ 0.32 cm) and the length of the narrow section was 0.5” (1.27 cm). The test specimen had a nominal thickness of 0.06” (~ 0.153 cm) but was measured for each specimen before testing. The test specimens were held with high-temperature pneumatic grips inside an oven (Instron Model # 3119.009 A2B1) set to 150 °C. The testing was controlled with a method developed using Bluehill 3 software (version: 3.66.4160). Video recordings of the tensile experiments provided supplemental information (e.g. shape, color, recoil, etc.) about the mechanical deformations to the break point at 150 °C. The test specimens were loaded into the pneumatic grips with a distance between the grips of 1.5”. They were allowed to equilibrate to 150 °C for 5 min. The specimen ends were then pulled/separated at a constant rate of 2.0”/min (5.08 cm/min) until reaching their tensile break point. The sample length L(t), and the applied tensile force F(t), was recorded at selected elongational times t. The data were collected at equal time intervals, and typically involved about 200 data points for each tensile study. The engineering stress (t) = F(t)/A0 at any time t was based on the tensile force F(t) measured at time t, and the cross-sectional area A0 of the tensile bar at t = 0 s. Similarly, the engineering tensile strain [(t) = L(t)/L0 – 1] was based on the sample length at time t and t = 0 s. Tensile properties of six to ten bars were studied for each material to provide estimates of the data precision.

FTIR Spectroscopy: A thin film of HPMC HME polymer was deposited onto the ATR crystal from chloroform (0.5 wt. % solution). Infrared spectra were acquired with a Thermo Scientific Nicolet iS50 FTIR and a heated Specac Golden Gate ATR accessory at a resolution of 4 cm-1. Forty-eight scans (71-second acquisition time) were collected for each spectrum. The ATR accessory was equipped with a single-bounce diamond ATR crystal. Spectra were collected from 50 to 250 °C (both heating and cooling) at a 10 °C/min



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rate. Peak heights at 3484 and 3596 cm-1 were measured relative to the baseline between 3200 and 3700 cm-1, and the peak height at 2929 cm-1 was measured relative to the baseline between 2700 and 3100 cm-1.

Size-exclusion chromatography (SEC): SEC sample concentrations were 2.0, 1.0 and 0.3 mg/mL for the samples having 15, 100 and 4M cP HPMC HME polymers, respectively, and were prepared in the SEC eluent (0.05 wt% NaN3 water solution). The solution was shaken at a flat-bed shaker at ambient temperature for two hours, and then stored in a refrigerator set at 4 °C overnight for complete hydration and dissolution. On the following day, the solution was shaken again for one hour, and filtered through a 0.45-µm nylon syringe filter prior to injection. Measurements were performed using a SEC system at 28 °C comprised of a Waters 2695 LC pump and autosampler equipped with Wyatt DAWN HELEO II multi-angle laser light scattering (MALLS) and Wyatt Optilab TrEX differential refractive index (DRI) detectors. The signals from the MALLS and DRI detectors were analyzed using ASTRA 6 software, version 6.1. Absolute molecular weight M and radius-of-gyration Rg were determined at each SEC elution volume increment using Debye plots (least-squares fit to a first-order polynomial according to the Zimm formalism) covering the light scatter angles between 25.8˚ and 132.2˚. A dn/dc of 0.140 mL/g for all HPMC samples was used.24 Reported molecular weight averages and radii of gyration were based on the first-order exponential fitting of slice M and Rg from ASTRA 6 software. Mass recovery was obtained by comparing the detected mass from integration of the DRI signal to the injected mass. The persistence length, lp, of the two higher-M materials was estimated from the SEC data (Rg vs M) based on the Kratky and Porod functional form.25-26 The fitting details and additional information are available in Supporting Information (Section S1). Table 1: Properties of HPMC HME materials and their aqueous solutions HPMC HME grade

Weight-average molecular weight (Mw , kg/mol) – powder1, 2

Dispersity (Đ)– powder1, 2

Wt.% MeO3

Wt.% HPO3

Radius of gyration (Rg,w , nm)1,2

Persisten ce length (lp, nm)4

15 cP

97.7

2.74

22-27%

25-32%

24.1

-

100 cP

204.8

2.56

22-27%

25-32%

35.7

4M cP

506.5

2.47

22-27%

25-32%

62.7

1.

Weight-average molecular weight (Mw , kg/mol)– plaque5

Tg (°C)6

97.9

100.3

6.2

200.0

98.7

8.3

476.3

100.0

As received. Average data of two aliquots reported.


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2. 3. 4. 5. 6.

From SEC data collected on aqueous solutions made with HPMC HME powders (as-is, without any processing) As determined by the United States Pharmacopeia and the National Formulary (USP-NF) substitution assay In aqueous solution. Calculated by fitting Kratky-Porod chain model to the SEC data (Rg vs. M). After compression molding. Average data of two aliquots reported. From DMS, sample warming rate 10 °C /min, measurement (angular oscillation) frequency 1 rad/s. Reported values are the peak temperatures extracted from the Gaussian functional from fitting to tan vs. temperature data on linear axes.

Results and discussions Temperature-dependent rheology of HPMC HME Polymers: The temperature dependent viscoelastic response under oscillatory shear (LVE regime) were studied for three grades of HPMC HME polymers at an angular oscillation frequency of 1 rad/s and the results are presented in Figure 1a-b. At low temperature, the G’ value was ~ (1.9  0.1) x109 Pa for all the three grades (i.e. insensitive to polymer Mw); this value is typical for amorphous polymers in their glassy state27 (Figure 1a). In this zone, G’ > G” signifying a solid-like material at 1 rad/s. Both, the G’ and G” decreased as temperature rose indicating a softening of the glassy materials as is commonly found for warming stable non-reacting materials not undergoing a thermal transition. Around ~ 90 -110 °C, a sharp decrease in both moduli was observed with warming, and was followed by an approach to a plateau modulus value [G’ ~ (2.3  1.1) x 106 Pa)] that is characteristic of amorphous high-Mw polymers or crosslinked rubbers in a rubbery state.27 Modulus measurements above ~ 150 °C were experimentally inaccessible due to loss of connectivity of the specimen between the top and bottom fixtures. The higher-Mw specimen (HMPC HME 4M cP) lost connectivity at a lower temperature ~ 120 °C. Several attempts were made with fresh aliquots of material and, in all cases, the specimens lost connectivity around the same temperature with some scatter (~  5-10 °C). The sharp decrease in moduli values between 90-120 °C was concomitant with a peak in the loss tangent (tan) as shown in Figure 1b. This thermal event was identified as the glass transition, and the peak temperature was extracted by fitting the data to a Gaussian functional form. The peak temperature or Tg were found to be 100.3  0.2, 98.7  0.4, and 100.0  0.2 °C for the HPMC HME 15, 100, and 4M cP 9 ACS Paragon Plus Environment


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polymers, respectively. Note here error values are the precision of the Gaussian fit to the data and not representative of replicate measurements ( 2 °C). The Tg values were insensitive of polymer Mw and consistent with previously reported Tg values measured using differential scanning calorimetry (DSC)20 but ~15 °C lower than those of broadband dielectric spectroscopy (BDS) Tg values.17 This discrepancy likely arises from a faster measurement rate (frequency) employed in the BDS (105 Hz or 6.3x105 rad/s) given the well-known dependence of the glass transition on the heating or cooling rate.28 A key observation was that all the three grades of HPMC HME exhibited solid-like behavior (i.e. the G’ > G” or tan < 1.0) at 1 rad/s even above their glass transition temperatures at applied stresses in the linear viscoelastic regime. The rubbery plateau (after glass transition) is commonly associated with either polymer chain entanglement or crosslinks. For typical thermosetting polymers (e.g. polyester, polyurethane etc.) this rubbery plateau continues (or even increase if the polymer cures) at elevated temperatures till the materials degrade or decompose. On the other hand, for typical thermoplastic semi-crystalline polymers (e.g. polyethylene), this rubbery plateau is followed by another sharp drop in moduli near the melting temperature. Selected semicrystalline polymers capable of hydrogen bonding (such as atactic polyvinyl alcohol, a long linear backbone polymer with OH side chains) crystallize when warmed near Tg (~ 190 °C) and therefore, a rise in modulus was observed between the Tg and onset of thermal degradation (~ 300 °C).29-30 However, HPMC HME are amorphous polymers10 where hydroxyl groups have been replaced by methyl ether and hydroxypropyl ether side chain grafting. Typical amorphous thermoplastic polymers (e.g. polystyrene) exhibit a substantial decrease in moduli as materials soften at elevated temperatures. But none of these usual classes of polymeric materials is known to demonstrate yield stress fluid-like rheology above their glass transition temperature (as will be described for the HPMC HME later in this work). Therefore, physically, even though the modulus value is consistent with a rubbery state, these HPMC materials did not reach a fluid-like state within our examined temperature range while oscillating at an angular frequency of 1 rad/s.


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(a)

(b) 10

HPMC HME

10

o

HPMC HME o

tan  = G''/G'

Pa

G',

10

or

8

7

10

10

HPMC grade 15 cP 100 cP 4M cP

6

5

Closed symbols: G' Open symbols: G"

-100

-50

o

Temperature sweep (-80 to 200 C, 10 C/min warming rate)

9

10

1.0

o

Temperature sweep (-80 to 200 C, 10 C/min warming rate)

10

G''

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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0.8 0.6 HPMC grade 15 cP 100 cP 4M cP

0.4 0.2 0.0

0

50

100

150

200

0

temperature T, °C

50

100

150

temperature T, °C

Figure 1: Temperature dependent (a) shear storage (G’, closed symbols) and loss (G”, open symbols) moduli; and (b) tan data for three different grades of HPMC HME polymers measured at an angular oscillation frequency of 1 rad/s. The solid red lines in Figure 1b are Gaussian fits to the data.

Prior to presentation of further results on HPMC, let us recall the temperature-dependent rheology of general-purpose polystyrene (GPPS), a common amorphous polymer having a similar glass transition temperature (see Supporting Information, Section S2). Linear oscillatory shear measurements show that well above the glass transition (Tg = 105 °C) the GPPS behaves as a solid (G’ > G”) at high frequencies; as a liquid (G’ < G”) at low frequencies; and the crossover frequency is temperature-dependent (ωcrossover ~ 0.04 rad/s at 140 °C, ~ 1.6 rad/s at 170 °C, and ~10 rad/s at 190 °C). At 170 and 190 °C, the empirical Cox-Merz rule31 was found to be valid (Supporting Information), signifying a rheologically simple system characteristic of most polymer melts and solutions. Shear-thinning behavior (in frequency space) can be described using the three-parameter Cross equation32 |𝜂 ∗ | =

𝜂0

(1)

[1+(𝜔𝜏)1−𝛽 ]

where |𝜂 ∗| is the magnitude component of the complex viscosity, 𝜂0 is the zero-shear viscosity, τ is a characteristic relaxation time, and β is the parameter that controls the steepness of viscosity reduction as a function of frequency. The best fit Cross equation parameters are summarized in Supporting Information, Table S1. These results serve as benchmarks describing a typical rheological behavior of a conventional entangled amorphous polymer at temperatures above its glass transition. 11 ACS Paragon Plus Environment

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Similar linear oscillatory shear experiments when performed on HPMC HME polymers at elevated temperatures, T = 130 – 230 °C, well above their glass transition (Tg ~ 100 °C based on the location of the loss tangent maximum, Figure 1b), the results were strikingly different. For all the three grades of HPMC HME polymers, data collected at two representative temperatures, namely at 150 and 230 °C, are shown in Figure 2a-d. At 150 °C, which is about 50 °C above the Tg, all the three grades of HPMC HME polymers demonstrated solid-like response (G’ > G”) in the entire range of experimentally accessible frequencies. A power-law fitting to the |η*| data returned an exponent value of ~ -0.8  0.06 when the same data are recast from the modulus format with G*= i* [where i = √(-1)]. Similar power-law frequency dependence of the | η*| was reported by Serajuddin and coworkers20 for these HPMC HME polymers for a temperature range 100 to 180 °C with a slightly lower scaling exponent value (~ -0.6). For these polymers, similar viscoelastic responses (G’ > G” in the entire frequency range and power-law scaling of complex viscosity), were observed for the temperature range in between 130 and 190 °C. (a)

(b)

HPMC HME o

T = 150 C

10

10

G''

or

G',

Pa

8

10

7

HPMC grades, modulus 15 cP, G' 15 cP, G" 100 cP, G' 100 cP, G" 4M cP, G' 4M cP, G"

6 Closed symbols: G' Open symbols: G"

0.1

1

10

100

1000

angular frequency , rad/s

(c)

10

(d)

6

HPMC HME o

5

10

or

G',

Pa

T = 230 C

G''

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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10

4

ω2 10

10

3

ω1

HPMC grades, modulus 15 cP, G' 15 cP, G" 100 cP, G' 100 cP, G" 4M cP, G' 4M cP, G"

2

0.01

0.1

1

10

100

1000

angular frequency , rad/s

Figure 2: Dependence of small-amplitude oscillatory shear flow viscoelastic data on the oscillation frequency for the grades of HPMC HME polymers at (a, b) 150 °C and (c, d) 230 °C. The frequency dependent shear storage (G’) and


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loss (G”) moduli data are shown in plots (a) and (c), magnitude component of the complex viscosity data are shown in plots (b) and (d). The dotted red lines in plot (b) and (d) represent power-law and Cross model fit to the complex viscosity data, respectively.

At elevated temperature (230 °C), a cross-over between the G’ and G” was observed for the HPMC HME low molecular weight (15 cP HPMC HME, Mw = 97.7 kg/mol) and high molecular weight (4M cP HPMC HME, Mw = 204.8 kg/mol) polymers (Figure 2c). For the intermediate molecular weight grade (100 cP HPMC HME, Mw = 506.5 kg/mol), an approach to crossover was observed at the lowest measurement frequency (ω = 0.1 rad/s). Note that while the crossover ω indicates a characteristic relaxation time, the typical liquid-like terminal rheology scaling (as ω → 0, G’ ~ ω2 and G” ~ ω1)27 associated with complete chain relaxation was not exhibited by any of these HPMC HME polymers. Instead, the system behaved like a cross-linked system close to gelation as evident from the similar frequency dependence of both moduli (G’ ~ G” ~ ω).33-34 Below in Table 2, we summarize the Cross model fitting to the complex viscosity data for all three polymer grades at 230 °C. For the cases where the low-shear viscosity plateau was not reached within experimentally accessible window, the estimated values of 𝜂0 should be treated as lower bounds. Note that the extracted 𝜂0 values followed a non-linear and even non-monotonic trend with the polymer chain Mw as the 100 cP polymer showed a higher value compared to the 15- and 4M cP polymers. This nonmonotonic trend likely not arising from different substitution levels as these three grades are chemically similar in nature (see Table 1) varying only in their molecular weights. A possible explanation of this observation is offered latter in this manuscript. Table 2. Cross model fit parameters to the oscillatory shear complex viscosity measurements at 230 °C for all the three HPMC HME polymers. For the 100 cP polymer the η0 and τ values are lower bounds, rather than exact estimates. HPMC HME grade 15 cP 100 cP 4M cP

η0 (Pa-s) (2.2  0.1) x 104 (40.6  2.6) x 104 (2.0  0.1) x 104

τ (s) 7.2  1.6 28.5  3.8 4.6  0.5

β 0.5  0.01 0.3  0.01 0.4  0.01

It can be argued that complete relaxation of these polymers could be accessed if the oscillation frequency range was extended to lower values (0.01 or 0.001 rad/s) or if independently probed with time-intensive creep and creep recovery studies. One practical limitation was long exposure of the polymer samples to 13 ACS Paragon Plus Environment

Biomacromolecules

elevated temperatures. A typical measurement sequence (including sample loading, strain amplitude sweep to determine the LVE regime and dynamic frequency sweep between the ω range 100-0.1 rad/s with 5 data points/decade) requires ~ 15-20 min. An extension of frequency sweep to another decade (ω value 0.1-0.01 rad/s) would have required an additional ~ 35 min of measurement time which might cause thermal degradation (e.g. caramelization, chain scission). The sample molecular weight distribution was measured by size-exclusion chromatography (SEC) before and after the rheological studies to assess any changes. The SEC-measured Mw data are presented in Figure 3. Compared to their unprocessed (powder) and compression-molded (plaque) values, minimal (< 10% decrease) changes in Mw were found for all the three grades of the HPMC HME polymers when exposed to a temperature of 180 °C or below for 15-20 minutes or less. However, at temperature of 190 °C or above, a significant reduction in Mw was observed likely arising from chain scission (Figure 3). Therefore, it was not possible (experimentally inaccessible) to measure zero-shear viscosity of these polymers with dynamic or creep studies. In addition, currently it is inconclusive whether the observed G’ and G” cross-over and an approach to complex viscosity plateau at 230 °C arise from polymer chain relaxation alone or from a convoluted effect of both the relaxation and thermal degradation. (b)

(c)

HPMC HME 15 cP

HPMC HME 100 cP

Processing conditions

600 500

200 100 0


After rheology at 230 C

w

300


400


0

M (kg/mol)





50

Plaque

100

Powder

150

w

M (kg/mol)

200


0

After extrusion at 190 C for 5 min.



Plaque


20

Powder

w

60


80

40

HPMC HME 4M cP

250

Powder

100

Plaque

(a)

M (kg/mol)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Figure 3: Weight-average molecular weight (Mw) data measured after different processing conditions. Note that the ordinate scales are different for these plots. During rheological measurements samples were typically exposed to elevated temperature for ~ 15-20 min. For the extruded sample the residence time was 5 min.


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At this point it is essential to comment that the typical residence time of polymer powders in hot melt extrusion process is 15s to 3min35 which is controlled by either screw speed (in single-screw extrusion) or feed rate/mass flow rate (in twin-screw extrusion). A longer residence, though possible and sometimes necessary for unit operations, is unfavorable as it may reduce line-speed and gives greater potential for thermal degradation of all formulation components. Also, the recommended upper temperature limit for HPMC HME processing is ~ 190 °C. The HPMC HME 15 cP polymer was extruded at 190 °C with a residence time inside the barrel of ~ 5 min (16 mm extruder). No significant Mw degradation was observed for the extruded specimen as shown in Figure 3(a). It confirms that the HPMC HME polymers are thermally stable under pharmaceutical-relevant extrusion conditions.

Yield stress fluid-like rheology In the foregoing section, it has been established that the HPMC HME polymers do not exhibit melt-like rheology under small-amplitude oscillatory shear at temperatures at least below 190 °C. Most conventional amorphous polymers begin to soften when warmed above their Tg and start to flow when warmed to temperatures about 30 to 50 °C above Tg. At sufficiently high temperatures (~ 70-90 °C above the Tg), meltlike rheology with complete chain relaxation is observed as elucidated for polystyrene thermoplastics (Figure S2, Supporting Information). Therefore, it is surprising that the HPMC HME polymers exhibit a solid-like behavior at temperatures as high as 190 °C, about 90 °C above their Tg. These polymers can certainly be extruded in the molten state well below 190 °C, or compression molded at 140 °C, as evident from the transparent strands and plaque photographs shown in Figure 4. These white friable HPMC powders can be transformed into a transparent cohesive strands in an extruder when the temperature exceeds the HPMC glass transition temperature, Tg. Similarly, these friable HPMC powders can be transformed into a transparent cohesive plaque with a compression molding process when the temperature exceeds the HPMC Tg and the applied compressive force is high (30,000 kg force). However, at the same temperature the plaques were found to be cohesive but opaque when the applied compressive force was lower. For both the extruded strands and plaques transparency and cohesiveness are common indicators of flow of the material in the molten state. 15 ACS Paragon Plus Environment


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Polymer viscosity generally decreases with increasing temperature and shear rate. One can argue that extrusion of HPMC HME polymers can be realized due to their shear thinning as these polymers experience high shear (as well as extensional) stresses and strains imparted by the flow path and screws in twin screw extrusion (TSE). In most relevant extrusion processes, shear strain rate in an extruder’s screw channels is about 10-300 s-1 and in dispersive mixing elements generally about 300-3000 s-1.36 However, these samples could also be compression molded which is essentially a zero- or negligible-shear process. Another possibility could be these polymers are thixotropic37 in nature where the flow is realized due to a change in microstructure (from gel-like to fluid-like) as a function of time at elevated (processing) temperature. Again the example of compression molding can be advanced as a counter argument where the processing (polymer flow followed by vitrification) time was ~ 15 min with limited (15 minutes). Therefore, the fluid-like behavior at elevated temperature (230 °C) likely arise from a combined effect of temperature (that resulted in polymer flow) and chain scission.


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Fluid-like (Severe MW degradation)

Tg

MW degrades at time > 15 minutes Fluid-like above a critical stress

100 110

130

150

170

190

Temperature ( o C)

210

230

Application extrusion temperature range under typical conditions Figure 7: Temperature-rheology-processibility behavior of the HPMC HME polymers.

Infrared Spectroscopy of HPMC HME: Hydrogen Bonding To gain insight into temperature-dependent physical state and stability of the polymers, infrared (IR) spectroscopy was performed on these materials. Temperature-dependence IR spectra of HPMC HME are shown in Figure 8a. The three major IR peaks were observed at ~ 1050, 2930 and 3490 cm-1. The strong peak at ~ 1050 cm-1 was assigned to dialkyl C-O-C stretch and the peak at ~ 2930 cm-1 was assigned to alkyl C-H stretches. The broad band at ~ 3490 cm-1 was assigned as the hydrogen-bonded OH stretch.44 The free OH stretch (i.e., non-hydrogen-bonded) was expected to be seen as a sharp band above ~ 3600 cm1

. A very small shoulder at 3600 cm-1 was observed at high temperatures as shown in magnified view in

Figure 8b. This band could be assigned as the free OH stretch, but in general, there was little change above and below the glass transition temperature, suggesting an unappreciable amount of hydrogen bond breakage. These results are consistent with a study45 that investigated hydrogen bonding in cellulose and found that its hydrogen bonds were stable up to 220 °C. The same types of temperature-dependent changes were observed in HPMC HME 15 and 4M cP polymers and data are presented in Supporting Information (Section S4, Figure S5). These spectral changes were reversible when the sample was cooled (Figure 8c), and there was no evidence of oxidation in the IR spectra after the heating cycle.



(a)

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(b) 50 C 87 C 123 C 159 C 196 C 232 C 250 C

C-O-C Stretch

50 C 87 C 123 C 159 C 196 C 232 C 250 C

CH Stretches

H-Bonded OH Increasing Temperature

CH Stretches OH Stretch

(c)

Heating: 50 C Heating: 123 C Heating: 196 C Cooling: 194 C Cooling: 121 C Cooling: 50 C

CH Stretches

OH Stretch

Figure 8: (a) Infrared spectra of HPMC HME 100 cP polymer from 50 to 250 °C. (b) Magnified view of the IR spectra for the same sample covering OH and CH stretches regions from 50 to 250 °C. (c) Infrared spectra of HPMC HME 100 cP collected during heating and cooling cycles.

The IR peak heights at ~ 3489 cm-1 (H-bonded OH stretch) and 3596 cm-1 (free OH stretch) were normalized to the CH3 stretch (2929 cm-1) to quantify the changes in the hydrogen-bonded and free hydroxyl groups as a function of temperature. The normalization accounts for changes in absorbance that are caused by changes in the contact of the HPMC film with the ATR crystal. The results for three polymer grades are shown in Figure 9. For the entire temperature window studied here (50-250 °C), the majority of the hydroxyl groups remain in the hydrogen bonded state. We note that this analysis considers both the intra- and intermolecular hydrogen bonding, but it does not account for changes in the peak shape of the OH stretch. The hydrogen-bonded fraction decreases steadily with temperature with a concomitant rise in free hydroxyl group fraction. This is consistent with the softening of the HPMC HME polymers at elevated temperatures.


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Another, interesting similarity to note that at any given temperature, the hydrogen-bonded OH fraction was similar for 15 and 100 cP grades and both were higher than that of the highest molecular weight (4M cP) grade. 0.5 0.45

OH Stretch/CH Stretch Ratio

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

0.4 0.35 0.3 0.25 0.2 0.15 0.1 0.05 0 50

100

150

200

250

Temperature (C) H-Bonded OH, 15 cP

Free OH, 15 cP

H-Bonded OH, 100 cP

Free OH, 100 cP

H-Bonded OH, 4M cP

Free OH, 4M cP

Figure 9: Temperature dependent free and hydrogen-bonded OH content for three different grades of HPMC HME polymers.

Potential Origins of Yield-Stress-Material-Like Behavior in HPMC HME above its Glass Transition To gain further insights into the temperature-dependent rheology of HPMC HME polymers, we simulated these

materials

using

the

Synthia

module

of

Materials

Studio

2016

(http://accelrys.com/products/collaborative-science/biovia-materials-studio/). We used a simplified HPMC repeat unit, with two methyl and one hydroxypropyl substitution per anhydroglucose unit (AGU). The estimated elastic modulus and zero-shear viscosity values for such an ideal HPMC polymer are presented in Figure 10. Those estimates are based on Quantitative Structure Property Relationships (QSPR) as outlined by Bicerano.46 The predicted Tg was 91 °C in a fair agreement with the experimental value (~ 100 23 ACS Paragon Plus Environment


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°C, DMS data). The temperature-dependent modulus predictions were in a good qualitative agreement with experimental data shown in Figure 1a. One notable observation was that well above the glass transition temperature (i.e. temperature above the sigmoidal change in the G’ data) a finite modulus value was found, in a reasonable agreement with experimental observations (Figure 10a). In comparison, for typical amorphous polymers such as polystyrene, predicted modulus values would drop to zero above the Tg as the materials transition into the melt state. Viscosity predictions, on the other hand, are more difficult to test, since zero-shear viscosities (associated with complete chain relaxation) were never attained within the experimentally achievable window (Figure 2a-d). One can only note that the predicted zero-shear viscosities for the 4M cP HPMC HME exceeded those found by extrapolating Cross-model fits of the data to zero angular frequency. Similar comparisons for the other two polymers (15- and 100 cP) at 230 °C were reasonably good. The yield stress was predicted to be around 80 MPa at room temperature (prediction not shown) and dropped to 0 MPa with warming as the temperature approaches the glass transition temperature (Tg ~ 90 °C). Based on QSPR analysis, one would expect HPMC polymers to behave as solid materials below ~ 100 °C and as viscoelastic fluids at elevated temperatures. It should be noted here that backbone hydroxyl groups in cellulosic materials are partially annihilated by the side-chain substitution process. This process is critical to break sufficient numbers of intra- and intermolecular hydrogen bonds, thus making the polymers more water-soluble. The HPMC HME polymers in aqueous solutions behave as semi-flexible chains as their calculated persistence length, lp, was found to be 8.3 and 6.2 nm for the 100 and 4M cP polymers, respectively (Table 1). These values are lower than the lp value for methylcellulose (MC) chains, ~ 13.3  0.3 nm, reported in literature.47 This is consistent with the fact that typically the side chain (methyl- or hydroxypropyl ether) substitution level per AGU is higher for the HPMC compared to MC. We believe that these substitutions break both the intra- and intermolecular hydrogen bonding that increases the backbone chain flexibility (mostly due to breaking of intramolecular hydrogen bonds) and at the same time lowers the glass transition temperature. From this argument one would arrive to a logical inference that the level of intramolecular hydrogen bonding in HPMC polymers


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is relatively low. However, intermolecular hydrogen bonding is still possible for HPMC as the hydroxyl group in the hydroxyl ether side chain can still form a hydrogen bond with oxygen atoms in the backbone. Based on this understanding we now aim to hypothesize a mechanism for temperature-dependent rheology of HPMC HME polymers. At temperatures below Tg, these cellulose ethers behave as conventional glassy polymers with elastic modulus G ~ 109 Pa, with contributions from intramolecular hydrogen bonds, as well as intermolecular van der Waals and hydrogen bonds. Once the polymer undergoes the glass transition, both van der Waals and intramolecular hydrogen bonds weaken or break altogether. However, intermolecular hydrogen bonds still remain and likely account for the modulus plateau at high temperatures. Applying the theory of rubber elasticity (G = RT/Mx where G is elastic modulus,  is density, R is universal gas constant, Mx is effective molecular weight between crosslinks and T is absolute temperature), we estimate the Mx value for these polymers to be approximately 0.4 – 0.5 kg/mol (using G values ~ 1 x 107 Pa for 15 and 100 cP or 8 x 106 Pa for 4M cP polymers from Figure 2a,  = 1100 kg/m3). If hydrogen bonds alone contribute to elasticity, then the calculations suggest there should be at least one effective hydrogen bond for about every two anhydroglucose repeat units. This network of intermolecular hydrogen bonds complements a more familiar entanglement network. Thus, the contributions to the HPMC modulus at high temperatures would come from both the entanglement elasticity and the network of intermolecular hydrogen bonds acting as temporary (albeit with fairly large lifetimes) crosslinks. (b)

(a) 1.0E+10 4M cP - model

4M cP - model

1.0E+14

100 cP - model

100 cP - model

1.0E+12

1.0E+08

4M cP - expt. 100 cP - expt.

15 cp - expt.

1.0E+06

, Pa-s

15 cP - model

G, Pa

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Biomacromolecules

15 cP - model 4M cP - expt.

1.0E+10

100 cP - expt. 15 cP - Expt.

1.0E+08 1.0E+06

1.0E+04

1.0E+04

0

100

200

300

0

T, C

100

200

300

T, C

Figure 10: QSPR predictions for (a) shear modulus and (b) zero-shear viscosity of HPMC polymer. Modulus values are predicted to be the same for all three molecular weights. Data points in Figure 9a are the same as in Figure 1a. Data points in Figure 9b are the η0 values reported in Table 2.



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The above hypothesis also helps explain why the dependence of modulus on molecular weight is nonmonotonic (Figure 2a and 2c). As the polymer molecular weight is increased, mobility of chains is reduced, and thus their ability to fully utilize all hydrogen bonding sites is hindered, albeit slightly. Thus, the degree of hydrogen bonding is somewhat lower for the 4M cP polymer, as compared to the other two grades (consistent with results reported in Figure 9). At T = 150 °C, the contributions of the hydrogen bond network is dominant, and thus G’ (15 cP) ~ G’(100 cP) > G’(4M cP).

At T = 230 °C, contributions from the

hydrogen bond network and from entanglement network become comparable. Thus, the 100 cP polymer has the highest modulus (strong entanglement and hydrogen bonding), while the 15 cP polymer (weak entanglement, strong hydrogen bonding) and the 4M cP polymer (strong entanglement, weak hydrogen bonding) have lower moduli. This hierarchy depends also on the frequency range considered in experiment, given that the hydrogen bonds are not permanent but have their own individual lifetimes, similar to socalled transient networks, such as associating thickener solutions48-52 or ionomers.53 The viscoelastic behavior of systems with multiple relaxation times are still poorly understood, with the fractional Maxwell model (FMM) approach showing some promise in describing the mathematical formalism,54-55 and the “sticky Rouse model” capturing many crucial physical details.53, 56 It should be noted that underlying reptation of entangled chains also influences the hydrogen bond relaxation time scales.56 Patashinski and co-workers57 recently proposed a new formalism to describe the yield stress behavior in systems with a distribution of effective bond energies; we imagine that a similar approach can be used to describe the yield stress of HPMC HME due to hydrogen bond network. Thus, the apparent yield stress would be a function of the temperature, the hydrogen bond energy distribution, the density of hydrogen bonds, and potentially even the applied external stress (see, e.g., ref. 58). Furthermore, the ductile/plastic (rather than brittle) nature of the polymer deformation at these temperatures seems to suggest that the “physical crosslinks” (most likely the hydrogen bonds) are constantly breaking and re-forming. One other remaining question is whether such an apparent yield-stress fluid like behavior is an exclusive feature of HPMC HME polymers or a generic characteristic of cellulose ether derivative polymer family. At this moment it is not possible to answer this question as most of the methyl- and hydroxypropyl 26 ACS Paragon Plus Environment

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methylcellulose polymers degrade before reaching their glass transition temperature. A separate and extensive undertaking of cellulose ether melt rheology is warranted in future to address this issue definitively.

Conclusions HPMC HME polymers exhibit yield-stress fluid-like rheology at temperatures exceeding its glass transition temperature. Only when the applied shear stress exceeds a critical value (σY) will these polymers flow like a fluid as needed for HME applications. The σY values under tension was measured to be ~ 0.5 MPa at 150 °C. These stress values are easily achieved in extruders used for hot melt extrusions of poorly-water-soluble APIs, while they are experimentally inaccessible using commercial rotational rheometers. An infrared spectroscopic study revealed that the majority of the hydroxyl groups in cellulose ether chains remain hydrogen-bonded (either intra- or inter-molecular) between 50 and 250 °C. We hypothesize that these hydrogen bonds act as physical crosslinks that do not relax unless the applied stress exceeds the critical stress value. These HPMC excipients, which flow as fluids only above a critical stress, may represent an exciting new class of thermoplastic materials.

ASSOCIATED CONTENT Supporting Information Experimental details on size exclusion chromatography (SEC), persistence length determination from SEC data, melt rheology of polystyrene, additional rheology and infrared spectroscopy data are available in the Supporting Information. This information is available free of charge via the Internet at http://pubs.acs.org/.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] (T.C.).

Notes: The authors declare no competing financial interest. 27 ACS Paragon Plus Environment


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ACKNOWLEDGEMENTS The authors acknowledge insightful discussions with Drs. William (Hunter) Woodward, Kurt Koppi and Mike Read. This research was supported by Dow Food, Pharma & Medical, a business unit of The Dow Chemical Company.

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Biomacromolecules

58. Eyring, H., Viscosity, plasticity, and diffusion as examples of absolute reaction rates. The Journal of Chemical Physics 1936, 4, 283.

Table of Contents (TOC) Graphic

Fluid-like (Severe MW degradation)

Tg

MW degrades at time > 15 minutes Fluid-like above a critical stress

100 110

130

150

170

190

Temperature ( o C)

210

230

Application extrusion temperature range under typical conditions


Rheology of Cellulose Ether Excipients Designed for Hot Melt Extrusion

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