Internal stress, molecular orientation, and distortion in injection moldings: Polypropylene and glass-fiber filled polypropylene

Internal Stress, Molecular Orientation, and Distortion in Injection Moldings: Polypropylene and Glass-Fiber Filled Polypropylene

C. S. HINDLE* a n d J. R. WHITE

Depar tment of Mechanical, Materials a n d Manufacturing Engineering

University of Newcastle Upon T y n e Newcastle u p o n T y n e , NEI 7RU United Kingdom

a n d

D. DAWSON a n d K. THOMAS

Division of Materials Metrology National Physical Laboratory

Teddington, Middlesex, TWI 1 OLW United Kingdom

Residual stress measurements and distortion analyses have been conducted on injection molded plaques made from polypropylene (PP) and a short glass-fiber filled polypropylene (GFPP). The residual stress analyses include measurements both parallel and perpendicular to the direction of flow during mold filling. Residual stresses are very anisotropic in GFPP, but not in PP. The residual stress levels in PP fall on aging at room temperature, whereas in GFPP the proportion of stress relaxation is smaller, and significant stresses remain even after heating to elevated temperatures. A significant contribution to distortion has been linked to the ejection process, and the long- and short-term distortion of moldings is discussed within the framework of the properties of the materials measured here.

INTRODUCTION esidual molding stresses and frozen-in molecu- R lar orientation both influence the properties of

injection moldings. Although quite separate charac- teristics, both are strongly dependent on molding conditions, in particular the thermal gradients that prevail during solidification. They are likely to be interactive, and it is difficult to separate them with confidence. We have measured residual stresses in straight bars by means of the layer removal technique (1 -4) in several studies of injection moldings (5-8) and have found that for noncrystallizing polymers the measured stress distributions in the as- molded state are in reasonable agreement with those predicted by simple theoretical studies [9, 101. Thus, the stress distribution through an injection molded bar or plaque resembles a parabola, with its peak (positive) at the bar center, indicating the location of

*Current address: Department of Chemistry, Napier College, Colinton Road, Edinburgh, EHlO 5DT United Kingdom.

the maximum tensile stress, and compressive stresses of rapidly increasing magnitude (up to approximately twice the maximum tensile stress) near to the surface, as in Fig. 1. This stress distribution is symmetrical if the cooling conditions within the mold are likewise, and therefore forces and mo- ments balance so that the bar or plaque remains flat. Post-molding conditioning such as annealing, especially in a temperature gradient, may change the distribution and in doing so may sometimes cause an imbalance to develop, which leads in turn to distortion (1 1).

The stress distributions for crystallizing polymers are generally more complicated than those found in noncrystallizing moldings. Aging at room temperature has been found to allow a reduction in residual stress level in crystallizing polymers with Tg below room temperature (6). Further experiments of this kind have been carried out and are described below, but the feature of particular importance in the work reported here is that a parallel investigation has

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("GFPP"-ICI HW60GR20) were molded using moving plate/fixed plate temperatures of 3O0C/3OoC, 8OoC/8O0C and 3O0C/8OoC. Melt temperature was 230°C. The changeover pressure during filling was 60 MPa for PP and 80 MPa for GFPP, and the con- trolled packing pressure was 28 MPa. All other set- tings (injection rate, cooling time, etc.) were kept the same for all batches of moldings. Each time the mold temperatures were altered, several moldings

I I Fig. 1 . Schematic diagram of a parabolic stress distribution plotted us. depth, z, in a bar. The positions of the surfaces are marked by the broken lines.

been carried out into the behavior of moldings produced in the same way but from a resin containing short glass fibers.

There have been many investigations into the effect of molecular orientation on mechanical properties (stiffness, strength, crazing behavior, etc.), but our concern in this work has been in post-molding distortion. The shrinkage of thermoplastics on molding (as distinct from that produced by post- molding annealing) is a well-known characteristic and one that is normally taken into account when a mold is being designed. The actual shrinkage obtained in a particular molding is, of course, dependent on the molding conditions, particularly on pressure up to the point of solidification. The imme- diate post-molding shrinkage can be measured by standard machine-shop tools, and warpage can also be fairly easily assessed (12). Further dimensional changes often occur during storage at room temperature (6). and the most successful method for moni- toring the rate of post-molding warpage changes has been the holographic technique (12, 131. In this paper, typical measurements of distortion are given and are related to the measured internal stresses.

EXPERIMENTAL

Specimen Preparation

Injection molded plaques were made in a tool having inserts that permit different runner systems, gating arrangements, and cavity geometry to be used, while maintaining accurate control of surface temperatures (14). They were molded in pairs using cavities 3 mm deep, one being a 90 mm square and the other a disc of diameter 100 mm. The volume of the square is 3% more than the disc, and the machine was usually operated with cavity pressure control using a pressure transducer in the square cavity. Polypropylene ("PP"-ICI GXM43) and glass-fiber filled polypropylene with 20 wt% glass

were rejected before collecting those to be used for testing. Plaques used for distortion measurement were allowed to age at ambient temperature, stand- ing on edge, and were used whole.

Moldings to be used for residual stress determination were allowed to cool for a time equal to ten machine cycles after ejection, and were then im- mersed and stored in liquid nitrogen until required.

C. S . Hindle. J . R. W h i t e , D. Dawson. a n d K . Thomas

I 58 POLYMER ENGINEERING AND SCIENCE, MID-FEBRUARY 7992, Vol. 32, NO. 3

Specimens were cut from the circular plaques, as shown in Fig. 2, and were again stored in liquid nitrogen if not to be investigated immediately. Spec- imens used to investigate aging were removed from liquid nitrogen and stored at room temperature for the chosen period before determination of their residual stress distributions. Some specimens were annealed in a recirculating air oven at temperatures of 80, 100, 110 and 130°C for periods of 1 h and 24 h before being cooled slowly to room temperature.

Measurement Methods

Determinat ion of Distortion

The most usual requirement in defining distortion is to measure out-of-plane displacement, which can be 0.1 to 1 mm. A convenient way of doing this is by formation of a moirt shadow fringe pattern (15). A line grating on glass, with spacing - 10 lines/mm, is placed over the object, and is illuminated from a white light line-source parallel to the grating lines at an angle of incidence 0 of say 45". Light reflected

B

G Fig. 2. The locations of bar-shaped specimens I , II, and III cut f rom disc moldings; G indicates the position of the gate.

lnternal Stress, Molecular Orientation, and Distortion

diffusely from the object surface escapes through the grating gaps. When viewed normally the object appears dark for those parts where distance d below the grating gaps is given by

d = (n+O.5)(linespacing)/tanB = O . l ( n + 0 . 5 ) mm

where n is an integer. Thus, in this example, a series of dark fringes is seen showing regions whose out-of-plane displacements differ in steps of 0.1 mm. The method shows the whole pattern of distortion in a very simple way. Once the pattern has been established, it may be sufficient to measure the height of the upper surface of the molding at particular points or along chosen lines with a dial gage.

Following the distortion that occurs during ejection (due to mechanical deformation) and during cooling (due to differential shrinkage and stress relief), further shape change due to unbalanced relief of stress continues for 2 to 3 days, after which very slow distortion may continue. Where required, this can be measured by two holographic techniques. The first is a sensitive double exposure laboratory method (12). The second is a real time method developed from a new simplified procedure of Ennos, et al. (16), and utilizes a single laser beam and a developed holographic image plate held close to a plaque (17); out-of-plane distortion of the plaque during aging causes one fringe to appear for every 0.3 pm displacement.

Determination of Internal Stress by Layer Removal

Thin layers were removed from straight bars cut from plaques using a single-point cutter on a high- speed milling machine. The resulting imbalance of forces in a bar containing residual stresses causes it to bend, and the curvature was measured after successive layer removals by a laser reflection technique described previously (6). Repeated removals and curvature measurements were made, and the analysis of Treuting and Read (1) was used to derive the stress profile, assuming in the first instance that the stiffness is the same at all levels within the bar. Although it has been found that a depth-dependent variation in stiffness of up to 10% (unfilled) and up to 40% (filled), dependent on direction, is in fact present in those moldings (18). we have shown that the influence on the derived stress distribution is fairly modest (19, 20) and that the unmodified distributions presented here are sufficiently accurate to make the comparisons described. The same com- ment applies to any variation of Poisson's ratio with direction or depth. Layer removals and curvature measurements were made as quickly as possible to minimize the effect of stress relaxation under the modified stress distribution, and the data for a single bar were collected in approximately 3 h.

Analysis Procedure

The Treuting and Read analysis (1) provides a formula that can be used to compute the residual

stresses in the xy plane at different levels, z l , within a sheet of material where the x and y axes are in the directions of principal stress and z1 is measured from the mid-plane:

In this equation, E is Young's modulus, and u is Poisson's ratio, both assumed uniform throughout the thickness and isotropic in the xy plane; and p x and p , are the curvatures measured in the x and y directions respectively. The upper and lower surfaces of the sheet are located at z = +. zo before layer removals commence. A similar expression for uy is obtained by interchanging subscripts x and y in Eq 1. For the studies reported here the x direction represents the bar axis direction, as in Fig. 3.

In the equibiaxial case, as should apply to a com- pression molded sheet allowed to cool in the absence of flow, it would be expected that ax = uy and p x = p , = p (say). Equation 1 then becomes (1,21):

1 20

+ 4 ( Z ~ f Z 1 ) P ( Z 1 ) - 2 1 p ( z l ) dzl (2) 2 1

There are two other important contractions of Eq 1. First, if the stresses are uniaxial, with a,,, = 0, then it is easily shown (21) that

Second, consider the case in which in one direction (y), the curvature is found to be very small after

'T

Fig. 3. Definition of the coordinate system. The bar, thickness 22,. has its axis parallel to the x-axis.

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160 POLYMER ENGINEERING AND SCIENCE, MID-FEBRUARY 1992, Vol. 32, NO. 3

each layer removal, i.e. p, = 0. Then Eq I becomes

It is notable that Eqs 2, 3, and 4 differ only in the factor outside the square bracket on the right-hand side.

In the case of a bar-shaped specimen, if x is chosen to coincide with the bar major axis, it is relatively easy to measure p x , but not at all easy to measure p,. If the bar is an end-gated injection molding it is to be expected that the stresses in the xy plane will be composed partly of those associated with thermal shrinkage and partly of those associated with molecules that were stretched out during flow. Computations of the residual stresses based on the thermal shrinkage mechanism have given reasonable agreement with measured stresses for unre- inforced thermoplastics (22, 231, and it is clear that a significant contribution must come from this source. If no other contribution were present then it would probably be valid to measure p x and to use it in the equibiaxial expression, Eq 2, making the implicit assumption that a,,, = q X and that p, , if it could be measured, would equal p x . If, on the other hand, there is a significant stretching and recoil effect, then this would cause q X to be of greater magnitude than ui,, and for p x to be greater than p,. In this case, if the measured value for px is placed in Eq 2 the result is an overestimate of stress, and if it is placed in Eq 4 the result is an underestimate. Since the difference in the two values will be a factor of (1 + u ) . which for most polymers is -1.4, the limits for ux are therefore reasonably well defined in this case, even without any measurement of p,. In our previous studies of bar-shaped injection moldings we have therefore not placed much emphasis on this point, and have merely stated our choice of formula.

In the work presented here we are concerned with plaques that have been molded using a gate located at the edge. In this case the thermal shrinkage and stretching stresses will add in a different manner at different locations. We therefore consider that the use of the biaxial form of the Treuting and Read analysis, Eq 1, is required. The ideal method of performing this analysis would be to remove layers from the whole plaque and measure curvatures in orthogonal directions at several locations. The mounting and machining of the whole plaque would be much more difficult than for bar-shaped specimens (although not impossible), and curvature measurement would be much more difficult and time-consuming. We have therefore chosen to cut bars from the plaques at different locations and in different directions and use these for layer removal experiments, using our usual method for curvature

measurement along the bar. Thus, one bar (I or I1 in Fig. 2). cut with its axis parallel to the flow direction, was used to obtain values of p X , and another bar (111 in Fig. 2) , cut with its axis perpendicular to the flow direction, was used to obtain values of p,. For convenience these values were used in E q 4 to calculate preliminary values u ; , ~ and u;,, for the residual stress distributions q X and ui.,. From Eqs 4 and 1 it follows that the corrected biaxial residual stress distributions are given by

and a,., = u;,, + uu;,x

As has been pointed out above, if the bar is cut from a plaque in the direction of flow and the curvature is measured in this direction the error in the computed stress in this direction will be no greater than a factor of (1 + u) . On the other hand, when the bar is cut in the direction across the flow so that the smaller curvature is measured, then the stress need no longer fall between the two values computed from Eqs 2 and 4, and the error is of unknown magnitude. This is why the biaxial form of the Treuting and Read expression has to be used. px has been obtained from measurements on bars cut parallel to the flow and p , from bars cut perpendicular to it. The results confirm that the stresses are quite anisotropic, and it seems almost certain that they will vary from one location to another in the plaques. Because the curvature measuring technique used here, in common with all of the other techniques applied to this kind of analysis, requires a bar several centimeters long, it has not been possible to perform measurements at several locations. If serious variations were present within the region from which the bars were cut, then the bar curvature would vary along its length. To test whether such variations in curvature were present, a series of circular arcs were drawn on paper, and at selected times during the layer removal experiments, the curved bar was placed against the arc of closest match for visual inspection. No variation in curvature was apparent in any of the examples inspected. The stress distribution results presented here can therefore be taken to be representative of a substantial region of the plaques. In passing, we consider that it is worth observing that this matching procedure could be developed into a most satisfactory curvature measurement technique.

In all layer removal experiments analyzed accord- ing to the Treuting and Read procedure, whether using Eq 1 or one of the modified forms ( E q s 2 or 4) . the dominant term in the stress expression when dealing with positions near to the surface is the first one, containing the gradient dpldz,. This value is extremely difficult to establish in this region because of the measurement interval (i.e., the depth of cut) and also because any systematic error arising from the introduction of stress by machining will

Internal Stress, Molecular Orientation, and Distortion

have its greatest effect here. Next to no significance can be attached to computed stress levels within the first interval of measurement, and it is advisable to treat all values within 0.2 mm of the surface with caution. A similar problem is encountered with GFPP near the bar center, because the change in curvature at each measurement interval is large, making measurement of the gradient of curvature somewhat imprecise. Although the relative importance of the gradient term is less here than near the surface, this remains a source of error. Once these limitations are recognized it is found that the re- producibility of residual stress plots from different bars with the same thermo-mechanical history is satisfactory.

RESULTS

Distortion

Examples of the moire shadow fringe results are shown in Fig. 4, which compares PP discs molded (i) conventionally with a small gate, and (ii) by cavity pressure control with a wider gate.

A convenient way of presenting the distortion results is to plot contours (derived from fringe measurements) of the top and bottom surfaces of the molding for selected sections. In the examples shown in Fig. 5, sections (i) parallel to the flow from the gate (GB), and (ii) across the flow through the center of the moldings (AA), are compared for both square and round GFPP plaques molded with both platens at 30°C. The shrinkage across diameter GB was approximately 0.9%, and that across AA was approximately 0.6% for both the square plaques and the discs. Less than 0.1 % variation in shrinkage was obtained with moldings produced with different mold temperatures.

Plaques were found to be distorted immediately after molding. Measurement shows that the distortion pattern is related to the ejector pin arrange- ment, and so comparative measurements must allow for this. For example, distortion will depend

(a) (b) Fig. 4 . MoirP shadow fringe patterns obtained with polypropylene: (a) molded conventionally (i.e., not with cavity pressure control) using a gate 1.5 x 1.5 mm; (b} molded with cavity pressure control, using a gate 3 x 1.5 mm.

F F

M c

mm 3- A

0 w

M F

mm

flow

Disc Square

Fig. 5. Summary of CFPP distortion measurements showing the location of both surfaces in perpendicular diametral sections for both a disc and a square plaque molded together. The diameters GB and AA correspond to the locations shown in Fig. 2. M indicates the surface molded in the moving plate, F the surface molded against the f i e d plate, with both surfaces at 30°C.

upon the degree of sticking to the mold, and upon the precise mode of operation of the ejector pins. With the present mold we have been able to make changes without disturbing the operations, and moirC fringe measurement on plaques that have stabilized at room temperature has shown:

with conventional processing, a decrease in filling speed by 10 x did not change the distortion; using cavity pressure control during filling gave moldings with much less distortion than those produced using preset machine conditions; with cavity pressure control, a 1.5 x 1.5 mm pin gate produced moldings that varied in weight by 2.5%. and the lighter moldings were more distorted; using a wider 3 x 1.5 mm gate reduced weight variation to 0.25%. and reproducibly gave moldings with slightly less distortion than the best smaller gate results; results for mold temperatures 3OoC/3O0C and 80°C/80"C were similar, but results for mold temperatures 3O0C/8O0C gave gross distortion; GFPP moldings were more distorted than those made with PP.

The holographic method has shown (12) that a molding largely stabilizes after 2 to 3 days at ambient temperature, and then, after it has stabilized for several months more, further out-of-plane distortion occurs at a rate that is just detectable (in one particular case for GFPP, 96% of relaxation bending occurred in 3 h after molding, and the bending rate fell by a factor of 30 in the next 5 h). Also, moirC measurement shows that annealing stabilized moldings for 1 h at 150°C produced negligible extra distortion.

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C. S . Hindle, J . R. W h i t e , D. D a w s o n , a n d K. Thomas

-10-

Residual Stress Analyses

Aniso tropy of S t res ses

Since the curvature across the bar width cannot be easily measured for bar-shaped end-gated moldings, it has not so far been possible to confirm the expectation that these moldings will have anisotropic stresses. In the case of plaques made from GFPP, the use of bars cut in different directions has confirmed strong anisotropy. Curvature versus depth-removed plots for bars cut parallel and perpendicular to the direction of flow in a round plaque molded with both platens held at 30"C, given in Figs. 6 and 7 , show significant differences. If, as a first approximation, the assumption is made for both cases that the curvature transverse to the bar axis is negligible, and E q 4 is used to obtain an assessment of the distribution of stress parallel to the bar axis, the results are shown in Figs. 6 and 7.

/

7

3

5

4 h

E

CL 1 v

3

2

1

0

Fig. 6. Layer remoual analysis for GFPP cut at location I! (parallel toflowjfrom a plaque molded with both platens held at 30°C: uncorrected stress distribution, u , ' , ~ . The double-headed vertical arrow marks the bar center.

This is the form we have previously used to repre- sent these results (24, 25). In Figs. 8 and 9 are shown the stress distributions obtained from the full biaxial analysis by application of E q s 5a and 5b. We believe that the results presented as Figs. 8 and 9 are more accurate than the unmodified plots shown in Figs. 6 and 7 , although the difference is not very great in the flow direction (x-direction). The bars cut transverse to the molecular flow

' O i

Fig. 7. Layer remoual analysis for GFPP cut at location 111 (across flow) f rom a plaque molded with both platens held at 30°C: uncorrected stress distribution, u ; , ~ .

Fig. 8. Corrected stress profile (solid line) f o r the jlow direction for the GFPP plaque molded with both platens held at 30". The uncorrected line from Fig. 6 is reproduced for comparison (broken line).

162 POLYMER ENGINEERING AND SCIENCE, MID-FEBRUARY 7992, YO/. 32, NO. 3


10

h

N E

r 2 5 Y

0

zo-z,(rnm)

Fig. 9. Corrected stress profile (solid line) for the transverse to flow direction for the GFPP plaque molded with both platens held at 30°C. The uncorrected line from Fig. 7 is reproduced for comparison (broken line).

showed considerable curvature about the bar axis after removing layers, and although this was quite clear on visual inspection, no measurements were attempted. At this stage we regard the use of two different bars to provide the p x and p, plots as the best practical measurement method.

The degree of anisotropy is much less marked with unfilled PP (Figs. 10 and 11). In this case, curvatures are so similar in both directions that a full biaxial analysis does not modify significantly the computed stress distributions (Fig. 12).

3

h N

. E 2 z I:

ti' v

. x

0.

-1

- 2

-3

j_io /[ c

9 PP 0.2 0.4 0.6 0-8 1.0 1.2 1.

zo-z (mm) 1

3

2

- E . v

Q

I

0

Fig. 10. Layer removal analysis f o r PP cut at location I1 f rom a plaque molded with both platens held at 30°C.

3

N-

I .-A

€ 2 1

b'

0

- 1

- 1

- .

0.2 0.4 0.8 1.0 1.2 1.4

3

2

E - Y

a 1

,O

Fig. 11. Layer removal analysis for PP cut at location I11 from a plaque molded with both platens held at 30°C.

Fig. 12. Results of a biaxial analysis based on the data in Figs. 10 and 1 I . u ; , ~ from Fig. 10 is reproduced ( . . . ' . ) for comparison with the corrected value ui,x (-.-.-.-). u ; , ~ from Fig. 1 1 is reproduced (-----) for comparison with the corrected value, ui,y (-).

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C. S . Hindle, J. R. W h i t e , D. Dawson , a n d K . Thomas

Effect of Mold T e m p e r a t u r e The residual stresses measured in plaques pro-

duced using mold temperatures as different as 30°C and 80°C showed remarkably little variation. Figure 13 shows results for PP molded with both platens at 80°C. A full biaxial treatment has been applied, and the stress distributions should be compared with the corresponding biaxial analyses presented in Fig. 12 for bars molded with both platens held at 30°C. There is perhaps closer agreement between the two directions of measurement at 80°C. Similar stress levels were obtained with plaques molded with one platen held at 80°C and the other at 30°C, although in this case greater initial curvature was obtained and was dealt with by the method described before (see Ref. 1 and below). It must be deduced that initial stress imbalances were rapidly removed by distortion immediately after ejection from the mold. Figure 14 shows results obtained with GFPP molded (a) with both platens at 80"C, and (b) and (c) with the fixed platen at 80°C and the moving platen at 30°C. The results are for the flow direction and are uncorrected for biaxial effects (i.e., u,' is plotted, not at), since it was established above that this does not produce substantial changes. The analysis presented in Fig. 14b is for a bar in which layers were removed from the fixed platen (80°C) side, and that in Fig. 14c is for layer removals from the moving platen ( 3 0 ° C ) side. The three stress profiles are remarkably similar.

Isayev and co-workers predicted that the level of

204

i

-2 -I /

'6

6

5

4

3

2 .. E

Q

-.. - 1

0

(b) Fig. 14. Layer removal analyses for GFPP cut at location I1 from plaques molded with (a) both platens held at 80°C, and (b) and (c) with fixed platen at 80°C. moving platen at 30°C. In (b), layers were removedfrom theflxed platen side: in (c) (next page), layers were removed from the moving platen side. In all cases the uncorrected stress

I1

-4 1 ' Fig. 13. Results of layer removal analyses for PP molded with both platens at 80°C. A biaxial treatment has been applied in both cases, and results are given for the depth dependence of stresses in the flow direction (-.-.-.-) and transverse to it (-). profile is presented.



(4 Fig. 14-Continued

residual stresses should not be very sensitive to mold temperature and confirmed this using measurements on polystyrene (26,271. On the other hand, Siegmann and co-workers found significant variations in residual stresses in moldings made from a modified poly(pheny1ene oxide) using different mold temperatures (28). Further references to cases in which mold temperature changes were found to al- ter residual stress levels are given in the paper by Siegmann, et al. (28).

Effect of Aging

The level of residual stress was found to fall on aging at room temperature with PP moldings, Fig. 15. This is similar to an earlier result obtained with a different grade of PP (6). Increasing the aging time from 6 to 12 months produced hardly any change in residual stress, Fig. 1%. A smaller reduction in stress levels was found with GFPP aged for 6 months under the same conditions, Fig. 16.

No difference could be detected in the residual stress levels in moldings stored in liquid nitrogen for prolonged periods. Figure 17 shows the curvature data points from Fig. 15b replotted for comparison with those obtained after liquid nitrogen storage (21 months in liquid nitrogen followed by 12 months aging at room temperature). The analyzed a, line in Fig. 17 is for the specimen stored in liquid nitrogen prior to room temperature aging. This confirmed an earlier result (6) and indicates that the storage of PP

I

- 1 I f Fig. 15. Residual stress analyses (no biaxial correction) for PP (a) aged 6 months at room temperature; (b) aged 12 months at room temperature (solid lines). Shown for comparison in (b) are the stress distributions obtained after 6 months ( . . . . . . ) and in the unaged bar (----).

specimens in liquid nitrogen is a suitable means of preserving the "as-molded'' state in a batch of specimens, which can then be tested in this state as time permits.

Effect of Annealing

In view of the modest stress levels found in PP and the appreciable reduction observed even with room temperature aging, it was decided to concentrate on GFPP for the annealing studies. In the previous sec- tion it was shown that room temperature aging had only a small effect on the residual stress levels in GFPP, but GFPP is often chosen for applications for which high service temperature capability is required. Hence, plaques were annealed for 1 h at temperatures in the range 80°C to 130°C. Our pre-

POLYMER ENGINEERING AND SCIENCE, MID-FEBRUARY 1992, VOI. 32, NO. 3 165

C. S . Hindle, J. R. White, D. Dawson, andK. Thomas

liminary results show that annealing at 80°C produces a small drop in stress levels, comparable with that observed after 6 months aging at room temperature (Fig. 16). A significant drop in stress levels occurs at 100°C. and still larger reductions were found at 110°C and 130°C. Increasing the annealing time to 24 h produces very little further change except at 80"C, for which further reduction in stress levels occurs.

An Example of Distortion on Aging

The first example of a distorted molding in which we investigated the residual stress distribution was a GFPP end-gated straight bar measuring 193 x 30 x 1.6 mm. It is common experience that thin flat moldings are prone to distortion. With this particular example the moldings were fairly flat and true on leaving the mold, but developed distortion within a few minutes. Measurements of distortion in these bars over a period of time have been reported previously (12). After many more months the distortion was in the form of a fairly uniform curvature at the

w- 20 E 5 z .bA 15 v

10

5

0

-5

-10

6

5

4

h

3 E . v

Q 2

1

0

Fig. 16. Residual stress analyses (no biaxial correctionl for GFPP (iJ unaged (-1 and (iiJ aged 6 months at room temperature (----).

O Q *, n.

0.8 1.0 1.2 1.4 z0-zJ rnrn )

-1 -1

2

A

E Q - v

1

0

-4 Fig. 17. Curvature plot obtained from a PP bar stored in liquid nitrogen for 21 months and then aged at room temperature for 12 months ( 0 ) . The corresponding stress proBle is shown. Data obtained with a similar bar aged for 12 months at room temperature, but not given a liquid nitrogen treatment, are shown for comparison (0); the stress proBle for these data is given in Fig. 15b.

gate end of the bar extending over almost one half of the length of the bar, while the other half of the bar remained almost straight, and it was at this time that the residual stress analysis was conducted.

The analyses for bars with an initial curvature were made using plots of curvature vs. depth of material removed in which the initial value of the curvature was subtracted from that measured at each successive layer removal, as prescribed by Treuting and Read (1). When this is done the stress distribution obtained is that which is present in the bar after it has reached the curved (distorted) shape in which it is in equilibrium in the absence of external forces. An alternative way of presenting the results would be to show the stress distribution in the bar if it were straightened by the application of external forces. Such a distribution could easily be derived from the initial curvature measurement, assuming uniform Young's modulus.

The results of layer removal analyses on bars cut from either end of the specimen that displayed end- to-end variation in curvature show that the overall level of residual stress is much smaller in the highly curved portion of the bar than in the part with smaller curvature (Fig. 18). This might as first seem to be contrary to the expected cause-effect relation- ship between residual stress and distortion. It must be emphasized that residual stresses do not produce distortion regardless of how large they are, provided they are in balance. That distortion was greater in the part of the bar containing smaller stresses sug- gests that in all likelihood the residual stresses were originally higher in this part of the bar (probably similar to those in the other half), and that stress relaxation occurred in a nonsymmetric manner, leading to an unbalanced distribution and change of

166 POLYMER ENGINEERING AND SCIENCE, MID-FEBRUARY 1992, VOl. 32, NO. 3


r

z-z (mm) 0 1 P’

(bl Fig. 18. Layer removal analyses for thin coupons (1.6 mm thick). (a) Curvature plots; (b) (uncorrected) residual stress plots for the bar axis direction. T h e solid lines correspond to the straighter end of the coupons; the broken lines correspond to the severely curved end.

ingly high distortion) if stress relaxation takes place nonsymmetrically .

DISCUSSION

The layer removal curvature plots obtained with GFPP show that the bar bends more than is the case with PP when this test is performed. The residual stress distributions derived from the curvature plots scale with the Young’s modulus, and so the difference in internal stress levels is greater than the difference in the curvature plots. The residual stress levels in GFPP injection moldings are then found to be much higher than those in unfilled PP, especially the tensile stresses in the interior-typically 20 MN/m2 (parallel to flow) or 7 MN/m2 (normal to flow), compared with 3 MN/m2.

One way of assessing the significance of the residual stress levels in bars made from different materials is to compare the residual stress level with the stress required to cause catastrophic failure, regardless of how this occurs. Thus with PP the ratio of the maximum tensile residual stress to the yield stress (> 30 MN/m2) is smaller than the ratio of the maximum tensile residual stress to the breaking strength (-60 MN/m2) with GFPP. Therefore it seems likely that the residual stresses will have a greater influence on failure in GFPP than in PP. In both materials significant compressive residual stresses were present near the surface, helping to protect the molding from failure by crack propagation from a surface flaw. This happens because freezing occurs inwards from the surface, and the later thermal contraction of the solidifying interior produces compressive stresses in the surface region, with balanc- ing tensile stresses in the center of the molding.

The structure of PP moldings is highly complex. The majority of molecules fold up into crystalline lamellae, which are themselves arranged into spherulites. The restricted space between lamellae and between spherulites contains other molecules, many of which have one or both ends incorporated into adjacent lamellae. The overall structure then consists of closely packed crystalline regions sepa- rated by amorphous material, which is to be re- garded as the continuous phase-but which is highly constrained by interaction with the crystalline regions. It is anticipated that any relaxation in response to residual stresses in the structure will mainly occur by rearrangement of molecular segments and movement of molecules in the amorphous phase, with accompanying elastic relaxation of stress in the crystalline regions, e.g., by change in molecular spacing. The morphology is known to vary through the thickness of a molding; and in particular for GFPP, the direction and degree of alignment of glass fibers can vary, with accompanying variation in stiffness and in thermal expansion coefficient.

In attempting to explain the differences in behavior between PP and GFPP, it is necessary to consider the possible sources of residual stresses in plastics

shape. The cause of such nonsymmetric relaxation cannot be traced, but this example does confirm that large residual stresses do not necessarily correspond to high distortion. On the other hand, if a molding is produced with large residual stresses, there is greater scope for the subsequent develop- ment of a significant imbalance (and correspond-

POLYMER ENGINEERING AND SCIENCE, MID-FEBRUARY 1992, Vol. 32, NO. 3 167

C. S . f f i n d l e , J . R. Whi t e , D. Dawson, a n d K . Thomas

additional to the conventional thermoelastic stresses that are produced in all materials (isotropic or other- wise) when solidification occurs in the presence of a temperature gradient. These are associated with the molecules that are not lying in their lowest energy configuration. This is typically caused during processing of the material into an article, when rapid deformation occurs with accompanying uncoiling and orientation of the molecules, and formation of a stressed state along and between molecules. The condition is then retained because complete relaxation of the molecules is relatively slow, especially if the material has been cooled after the forming operation. The material may in fact be cooled below its glass transition temperature, T g . where only very local bond movements are possible and any stress present could have only a small influence in varying molecular conformations (27, 29, 30). However, with polypropylene at room temperature, the material is some 20°C above Tg, and the typical bond relaxation time is only s for “glass-rubber” conformational changes in amorphous regions (the relaxation for polypropylene). Thus the initial expectation is that considerable movement of molecules in the amorphous phase within and between spherulites could occur in response to the stressed state. The molecules forming crystalline lamellae in the spherulitic structure are believed to be largely immobilized in their lattice at temperatures up to, say, 80“C, but the a-relaxation process found with PP is ascribed to very slow translations of chains within the lamellae, which yield coupled rearrange- ments of amorphous segments (31), and could con- tribute to slow relaxation of stress. In fact it is found that shape change occurs for several days after injection molding (e.g., 12). and that internal stress declines (6). Nevertheless it is here shown that after one year at room temperature, appreciable stresses remain in PP (Fig. 17). To account for this, it is considered that some amorphous phase molecules are effectively pinned at points where they interact with-and very likely become part of-the crystal lamellae.

I t is not possible to say how much higher than the stresses measured here are the stresses induced during melt processing, for much relaxation must occur during cooling. It is often observed that a molding can distort rapidly after ejection from a mold, under the action of internal stresses-indeed, it is common practice to “jig” the molding to hold the shape until the material is sufficiently frozen.

Isayev and Crouthamel have observed anisotropic residual stresses in a glassy polymer (polystyrene) and have discussed their origins in some detail (27). In the present work the polymer is semicrystalline with Tg below room temperature, leading to differences in behavior. To explain the presence of anisotropic stresses in the plaques studied here we suggest two principal mechanisms: (i) Recoil of molecules stretched out during flow

will cause shrinkage in the same direction as the

(ii)

molecular orientation (this is the counterpart of die swell in extrusion processes). The degree of molecular orientation differs with position in a melt flowing into a cold mold, being high adjacent to the solidified layer on the mold surface. There is also a variation in the cooling rate through the molding, so that greater relaxation of oriented molecules occurs towards the molding center. The combination of both effects may give rise to differences in extension and orientation of molecules at different depths before the material freezes, and hence to residual stresses; and anisotropic molecular orientation would then result in anisotropic stresses. These entropy elastic stresses are actiue (32, 33). just as energy elastic stresses are, whereas in glassy polymers they are pass iue and will have very little influence on the layer removal results. Following freezing, thermal contraction occurs during further cooling. If the material retains preferred orientation at the point of solidification, the thermal contraction may be anisotropic, producing further anisotropy of residual stresses.

Thermoelastic calculations of stress profiles that take account of these factors have not yet been published to our knowledge, but a significant effect may reasonably be expected. The failure to observe stress anisotropy with the present PP moldings can be attributed both to restricted orientation of molecules on filling this mold, and to rapid relaxation of molecular orientation at temperatures too high to induce residual stress. Studies of the morphology of injection moldings (34-40) have shown that except for a surface layer of <0.6 mm thickness, the material forms into nearly equiaxed spherulites and is thus expected to be nearly isotropic. The absence of anisotropy of residual stress in the plane of the molding then leads to the conclusion that a large fraction of the measured stress comes from the volumetric shrinkage men- tioned above.

On the other hand, larger and anisotropic residual stresses are found for GFPP (e.g., Figs. 6 and 7). With this material the glass fibers are arranged anisotropically in the solidified molding (41 -44). and the structure is commonly stratified with regions of quite different orientation distributions. At present we do not have any means of assessing quantita- tively the effect of fiber orientation on residual stress values, and so it would serve no purpose to consider fiber orientation distribution in detail. However, it is relevant to note that:

(i) The fibers produce higher overall moduli than for PP (3.7 GN/m2 cf 1.2 GN/m2 in the flow direction). Thus any internal strains arising from the foregoing mechanisms would develop higher residual stress (for the same reason that stress calculated from curvature developed on layer removal scales with modulus).

(ii) In a similar manner, the anisotropy of moduli


ln ternal S t ress , Molecular Orientation, a n d Distortion

arising from fiber orientation is expected to produce anisotropy of internal stress.

(iii) The presence of higher moduli and greater anisotropy shows that higher anisotropic con- straints exist on the movement of molecules at room temperature. The glass fibers, coated to promote interaction with polypropylene, will to some extent provide pinning points for molecules. However, the more significant effect probably derives from the structure formed on mold filling. The presence of the fibers will in- hibit relaxation prior to solidification; and they also act as crystal nuclei, causing “transcrystal- line” growth normal to the fibers (45-47). Thus in the volume of material around fibers it is expected that more molecules will be retained in a strained condition, and trapped in this condition in the amorphous regions. The conse- quence would be higher and anisotropic residual stress.

(iv) The presence of the glass fibers will also affect the thermal contraction on cooling below the freezing point, again causing higher internal stresses and greater anisotropy to develop. Stresses could be further increased by the variation of fiber orientation through the thickness.

(v) Lastly, it should be observed that analysis of layer removal measurements so far developed does not take into account the effect of having anisotropy of Young’s modulus. However, it is felt that the influence will be modest, as when calculating the effect of a variation of modulus with depth (see above).

All these effects are additional to those occurring in PP (which will occur also in GFPP) and provide credence for the experimental finding that residual stress is anisotropic in GFPP, and higher than in PP.

The results of the aging tests can now be assessed, for they are consistent with the ideas advanced to explain the anisotropy. The fractional relaxation of residual stress was observed to be much less with GFPP than with PP. No relaxation is anticipated in the glass phase, but since this occupies only -10% by volume, this cannot explain the differences in behavior. Instead we reiterate the hypothesis that the GFPP moldings contain a significant number of molecular segments that are extended in the direction of flow and pinned at both ends. We envisage pinning, such as that provided at a fiber or at the entry of a molecule into a crystal, to be much more permanent than a simple entanglement, so that the conformational (recoil) contribution to residual stress is unlikely to be relaxed easily, thus account- ing for the retention of a high fraction of the residual stress magnitudes on aging. In the case of unfilled PP we deduce from the low level of stress that less pinning is initially present. In both materials some molecular relaxation will inevitably occur between points of entanglement or pinning. However, the annealing experiments show that quite high tem-

peratures are required ( 2 100°C) before substantial changes occur with GFPP. This material has a good reputation for applications demanding a high service temperature, and greater structural stability provided by molecular pinning could be the reason.

The residual stress levels were found to be sur- prisingly insensitive to the mold temperature. When the two mold platens were at different temperatures, this caused the molding to warp, but the residual stresses in the distorted configuration were not very different from the symmetrical ones obtained for the specimens molded with a uniform mold temperature. It is recalled that the method of analysis used here provides the stress distribution through the molding free from external stresses. Thus, it is per- fectly acceptable for a warped molding to have a symmetrical or nearly symmetrical stress distribution. If, however, the residual stress distribution were to be presented for the molding when constrained to be flat, this would demand the application of an external bending moment, and the bending of the molding to restore flatness would create within it a nonsymmetrical stress distribution.

With GFPP the residual stress levels obtained when both platens were held at 30°C were greater than those when they were held at 80°C. This is consistent with simple thermoelastic analysis, which predicts that the residual stress levels should be greatest when the temperature difference between the melt and the mold cavity wall is greatest. Plaques molded with the fixed platen at 80°C and the moving platen at 30°C gave residual stress levels inter- mediate between those obtained when both platens were at 80°C and 30°C. respectively, with the side of the plaque adjacent to the 30°C wall giving marginally higher stresses than that adjacent to the 80°C wall. Since stress distributions are presented for the plaques in the deformed state, free from external loading, we expect that the bending mo- ments derived from results like those in Figs. 14b and 14c-obtained by layer removal from 80°C and 30°C sides, respectively-should balance one another. This is confirmed by the results in Figs. 19a and 19b, in which zlu, is plotted vs. z1 using the data in Figs. 14b and 14c, respectively. The areas of these plots correspond to 1 z lui d z , the bending moment about the center, and in this example the areas differ by approximately 5%. (Note that in both cases the area is negative: the numerical value of the positive part in the low z1 region must be subtracted from the numerical value of the negative part in the high z, region to produce the required integral.) This agreement can be considered to be good in view of the admitted inaccuracy of layer removal residual stress measurements near the bar surface, and because we have not attempted to cor- rect the analysis to account for variations in Young’s modulus through the bar thickness (19, 20).

If subsequently, on aging or annealing, stress relaxation should take place unequally in the two halves of the above plaque (as might result from

POLYMER ENGINEERING AND SCIENCE, MID-FEBRUARY 1992, YO/. 32, NO. 3 169

C. S . Hindle, J . R. Whi t e , D. Dawson, a n d K . T h o m a s

5 1

\ t I I

I

- 1 5 i

5-1

z,(rnrn)

0 pof, , , , P O , , , ,'"

I \ I

-1 5 1 (b)

Fig. 19. Plots of Z ~ C J ; , ~ versus z I . (a) Constructed using t h e d a t a i n Fig. 14b; (b) constructed using t h e d a t a i n Fig. 14c.

different stress levels or from morphological differences), then the shape of the molding would change.

We are now in a position to consider distortion. In common with experience in the field, GFPP moldings tended to possess more distortion on removal from the mold than those made from PP. The higher residual stress may be responsible for this, although only the imbalanced component can actually cause distortion. Curvature caused by nonsymmetrical residual stresses can be expected in plaques molded with the platens at different temperatures, but other causes must be sought to explain distortion in plaques molded with the platens held at the same temperature. The observations show a strong effect connected with ejection. If the molding resists ejection and a high force has to be applied, causing it to deform during the process of ejection, then in the case of a glass-fiber filled polymer the deformation may be sufficiently severe to produce permanent relative displacements of the fibers within the molding, leaving behind a permanent deformation. This

is especially likely if regions in the interior of the molding have still not solidified at the point of ejection. An equivalent deformation in PP should re- cover more completely: ejection forces with plaques and similar shapes are likely to be smaller in any case, because of the greater shrinkage of PP. One other likely cause of higher distortion with GFPP is anisotropic shrinkage arising from glass fiber orientation.

CONCLUSION

Residual stresses have been found to be highly anisotropic in plaques molded from GFPP, whereas residual stress anisotropy is not pronounced in those molded from PP. Residual stress and molecular and glass fiber orientation all influence distortion in injection moldings. In newly molded parts, forces applied by the ejection system may cause distortion, although some of this distortion may be recovered. Recovery is less likely with filled polymer like GFPP. Residual stress magnitudes decay on aging at room temperature or at elevated temperature. If the stress relaxation has an out-of-balance component, then a contribution to distortion will develop. Even though higher residual stress values are observed in GFPP than in PP, they decay less readily at room temperature so that, once the part is molded, it is likely to have better dimensional stability if made from the reinforced material. This is true also at elevated temperature. Thus, if the molding conditions, in- cluding the ejection arrangements, can be adjusted to produce parts from GFPP with good dimensional accuracy, they are more likely to stay within specifi- cation than if they are made from unfilled PP. The same is likely to be true for other semicrystalline polymers that are used both with and without reinforcement.

ACKNOWLEDGMENT

The authors are grateful to Dr. B. E. Read of the National Physical Laboratory (United Kingdom) for discussion of internal stress and relaxation mechanisms; and financial support from the SERC is gratefully acknowledged.

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Documents

Internal stress, molecular orientation, and distortion in injection moldings: Polypropylene and glass-fiber filled polypropylene