Massive stress changes in plasma-enhanced chemical vapor deposited silicon nitride films on thermal cycling

Thin Solid Films 460 (2004) 7–16

Massive stress changes in plasma-enhanced chemical vapor deposited silicon nitride films on thermal cycling Michael P. Hughey*, Robert F. Cook Department of Chemical Engineering and Materials Science, University of Minnesota, 421 Washington Ave SE, Minneapolis, MN 55455, USA Received 10 July 2003; received in revised form 26 November 2003; accepted 7 January 2004 Available Online 24 March 2004

Abstract Massive irreversible increases in tensile stress (up to 2 GPa) on thermal cycling are demonstrated for plasma-enhanced chemical vapor deposited (PECVD) silicon nitride films. Results give further evidence for the claim that this phenomenon is generic to PECVD films and is attributable to the removal of bonded hydrogen: the magnitude of stress increase is independent of the film stress and can be accounted for with a calculation involving the amount of evolved hydrogen. The massive stress changes cause film fracture in most of the films discussed here, with a large diversity of fracture behavior exhibited. The effects of deposition conditions (temperature, plasma frequency, substrate) on film modulus, hardness and coefficient of thermal expansion, as well as stress and stress hysteresis are also examined. 䊚 2004 Elsevier B.V. All rights reserved. PACS: 68.60.Bs; 68.60.Dv; 62.20.M Keywords: Stress; Silicon nitride; Heat treatment; Chemical vapor deposition

1. Introduction Plasma-enhanced chemical vapor deposited (PECVD) dielectric films are ubiquitous in the microelectronics industry as well as in other advanced technologies. As fabrication of most microelectronic devices requires many thermal cycles for the deposition or annealing of each layer of material deposited, full understanding of the thermo-mechanical properties of PECVD films is required to help select materials and processing steps to optimize mechanical reliability of devices, as stress development during fabrication is unavoidable. The measurement and prediction of the development of reversible stress caused by thermal expansion mismatches between device materials is well understood. However, it is also well known that various thin film materials can exhibit hysteretic or irreversible stress changes on thermal cycling. These latter phenomena are *Corresponding author. Tel.: q1-612-6267410; fax: q1-6126267246. E-mail address: [email protected] (M.P. Hughey).

consequences of the non-equilibrium nature of thin film deposition processes, the mechanical constraint of the substrate, or both. Metal films exhibit hysteretic stress responses on thermal cycling as stress that develops due to a thermal mismatch between the film and substrate is relaxed by creep deformation processes at high temperatures w1– 5x. A typical stress-temperature response includes compressive thermal stress development on heating, followed by relaxation to near zero stress at high temperatures, followed by tensile thermal stress development on cooling. Attempts at modeling or predicting such stress development by utilizing bulk creep deformation models have been relatively successful w1–3x, although the predictions do not describe all the differences between thin film and bulk mechanical behavior w2,3x. Less well known than the stress behavior of metals are the irreversible stress changes that occur in dielectric films deposited by non-equilibrium processes that are subjected to thermal cycling or annealing. This phenomenon is primarily observed in PECVD dielectrics containing hydrogen: SiOx:Hy w6,7x, SiNx:Hy w8,9x,

0040-6090/04/$ - see front matter 䊚 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2004.01.047

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M.P. Hughey, R.F. Cook / Thin Solid Films 460 (2004) 7–16

SiCx:Hy w9,10x and SiOxNy:Hz w11,12x, for example. However, irreversible stress development has also been observed in other dielectrics such as r.f. sputtered aluminum oxide w13x, r.f. magnetron sputtered silicon carbide w14x, and silicon oxide and aluminum oxide deposited by reactive electron-beam evaporation, plasma ion-assisted deposition, and ion-beam sputtering w15x. Unlike in metal films, the hysteretic stress responses in these films do not appear to be driven by the relaxation of reversible thermal stress changes. Bulk ceramic counterparts of these materials deform elastically and by brittle fracture—they do not have active creep deformation mechanisms at the moderate temperatures at which films are frequently tested (below 700 8C). It is commonly accepted that irreversible stress develops in PECVD films from the reduction of the amount of bonded hydrogen in the film, but it is unclear exactly how this mechanism occurs. Furthermore, other explanations have been offered, including the release of unbonded molecular of hydrogen w8x and void shrinkage w16x. The lack of an accepted explanation for irreversible stress development has hindered the formation of physically based quantitative models of stress-temperature behavior for hydrogen-containing dielectric films analogous to those for metal films. In this paper, silicon nitride is used for study as a typical PECVD film. Silicon nitride is advantageous in that it is not as susceptible to moisture absorption as silicon oxide, it is chemically simpler than silicon oxynitride, and its stress behavior is more frequently reported in the literature than that of silicon carbide. Technologically, silicon nitride is important as a passivation layer in microelectronic devices, as an etch stop layer in interconnect stacks, and as an active element in optical waveguides. In particular, production of optical waveguides can require high temperature thermal cycling and annealing to reduce the amount of incorporated hydrogen in the PECVD materials w17x, requiring mechanical stability of the dielectric to prevent deleterious consequences—primarily shape change induced by stress, and fracture. The goals of this paper are, then, to (1) identify the general features of irreversible stress development in PECVD films to establish a framework for subsequent quantitative modeling and prediction, which could then assist in design choices in a variety of technologies employing PECVD films, and (2) identify specifically the deposition conditions suitable to maintain mechanical integrity of a silicon nitride film in optical devices. 2. Experimental procedure Eight silicon nitride films, roughly 200 nm thick, were deposited by PECVD by varying the temperature — 320 or 120 8C—the plasma frequency — 13.56 or 0.46 MHz—and the substrate — InP or Si. The single

crystal substrates were (100) oriented, 50 mm in diameter, and either 385 mm thick (InP) or 300 mm thick (Si). Gas flows were 1000 sccm 2% SiH4 in N2 and 12 sccm NH3; the pressure was 600 mT, and the r.f. power density was 61 mWycm2. Film thickness was measured by ellipsometry with a helium–neon laser (before and after thermal cycling) and confirmed by cross-sectional scanning electron microscopy (after thermal cycling). Wafer curvature measurements were performed using an FSM900TC (Frontier Semiconductor Inc.) and film stress was calculated using the Stoney equation w18x for biaxial film stress, sf: sfsŽkyk0.

E¯ st2s , 6tf

(1)

where k0 and k are wafer curvatures before and after film deposition, E¯ ssEs y(1yns) is the substrate biaxial modulus (modulus, E, and Poisson’s ratio, n), and t is thickness with subscripts appropriate the film, f and substrate, s. The values used for substrate biaxial moduli were 180.5 GPa for Si w19x and 94.9 GPa for InP (calculated w20x from the elastic constants w21x). The films were thermally cycled at a rate of 2 8Cymin to a peak temperature and back to room temperature; peak temperatures were sequentially increased to a maximum of 630 8C. This was performed under a nitrogen atmosphere vented to a fume hood due to concerns of InP dissociation. At each temperature, two values of curvature were measured for each of four orientations (spaced by 458). The two measurements for each orientation were averaged, and the means and standard deviations of these four values were taken as the average and uncertainty in film stress. (Here, uncertainty refers to the error one assumes in calculating a single value of biaxial film stress for a non-equibiaxial stress.) The curvatures of the bare substrates were not measured beforehand, and stress was estimated by assuming the bare substrates were perfectly flat (k0s0). The uncertainty introduced by such an assumption is at most 120 MPa for films on InP and 70 MPa for films on Si, based on unpublished experience in this laboratory for the largest possible bare substrate curvature. Although these uncertainties are large, the magnitudes of irreversible stress changes discussed in this paper are much larger, and precise knowledge of the film stress is not necessary to understand the stress change phenomenon. The plane-strain modulus, E*f sEfyŽ1ynf2., and hardness, Hf, of the films were measured by depth-sensing indentation using a Nano Indenter䉸 XP (MTS Systems Corp.). However, these measurements could provide only an estimate of the true values due to the thinness of the films and the hard film-on-soft substrate problem w22x (particularly in the case of InP). A Berkovich diamond tip that was well calibrated on fused silica to a minimum contact depth of 10 nm was used; continuous

M.P. Hughey, R.F. Cook / Thin Solid Films 460 (2004) 7–16

stiffness measurements of E*f and Hf were averaged over the range from roughly this value to approximately 10% of the film thickness, over which E*f and Hf were relatively constant. The coefficient of thermal expansion (CTE), af, of the films was calculated by using the slope of the stress-temperature data (near room temperature where stress develops due only to a difference between af and as) and E¯ s in the thin film approximation: 1 dsf afsasy ¯ Es dT

(2)

where T is the temperature. The substrate coefficients of thermal expansion were obtained from the literature, and the room temperature values used were 2.6 ppm Ky1 for Si w23x and 4.56 ppm Ky1 for InP w24x. For calculation of biaxial modulus from plane-strain modulus, Poisson’s ratio was assumed to be 0.2—typical of an amorphous ceramic (and slightly less than that of bulk polycrystalline silicon nitride w25x). Use of the socalled ‘double substrate method’ w26x to calculate af and the film biaxial modulus, E¯ f, was precluded because of obvious differences in the measured properties of films deposited under the same conditions on different substrates. Finally, the hydrogen and total atomic concentrations were measured by forward recoil spectrometry (FReS) and Rutherford backscattering spectrometry (RBS) for some films. RBS experiments were conducted using 3.4 MeV Heqq ions, a backscattering angle of 1658, and a total charge collection of 20 mC (40 mC for films deposited on InP). FReS spectra were acquired using 3.0 MeV Heqq ions incident at 158; the detector was 308 from the beam direction, a 12 mm thick Mylar䉸 foil was used to stop scattered He ions, and 20 mC of charge were collected. The hydrogen concentration was calculated by using a Kapton䉸 (polyimide film, C22H10O5N2) standard. Hydrogen and total atomic concentrations were calculated with approximately 5–10% relative uncertainty. 3. Results and discussion 3.1. Film stress and mechanical properties The estimated film deposition stress—that is, the intrinsic stress plus thermal stress developed on cooling from deposition temperature to room temperature—is shown along with the film modulus, hardness, and CTE in Table 1. Several trends were observed in the film properties. While keeping substrate and plasma frequency constant, hardness and modulus always increased with increasing deposition temperature, although the increase in modulus was not always statistically significant. With one exception, modulus and hardness also

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increased with decreasing plasma frequency for a given substrate. The CTE decreased with increasing deposition temperature for a given substrate, and was invariant with frequency on Si for a given temperature, but increased with increasing frequency on InP (regardless of temperature). The scaling of modulus and hardness with temperature and plasma frequency matched previous observations w11,27x and is explained by greater film mass density. For every film, the modulus was smaller than for bulk silicon nitride (308–318 GPa w25,28x), but was nonetheless large for a PECVD silicon nitride film; the hardness values were in the range observed for bulk silicon nitride (14–25 GPa w28,29x). The CTE of half of the films was approximately 3 ppm Ky1, the value for bulk polycrystalline Si3N4 w30x, whereas the remainder of the films had a smaller CTE. Not much change is expected in the modulus or hardness on thermal cycling, although large changes are expected for the CTE w6x. Finally, although it is not clear why the CTE is apparently affected by the substrate, the smaller values of modulus and hardness for films on InP is most likely attributable to a decrease in apparent film properties due to the softer, less stiff substrate’s influence on the measurement. (Incidentally, if the same modulus value were used to calculate the CTE of two films deposited in the same way but on different substrates, the calculated CTE’s of the two films would be even more different than as shown in Table 1.) Different trends were observed for the deposition stress: films were more compressive (negative stress) for lower frequency plasmas, as is expected w27,31x due to greater ion bombardment under these conditions w32x; no dependence on deposition temperature was observed; and films were always more compressive on Si. The lack of dependence of deposition stress on deposition temperature is not necessarily unexpected: although stress typically increases with increasing deposition temperature w33x, this is not always observed w11,34x. That films on different substrates should have different deposition stresses is expected due to a difference in thermal stress; however, the calculated thermal stress developed on cooling from the deposition temperature to room temperature is more compressive for a film on InP vs. a film on Si (by only 36 MPa and 106 MPa for 120 8C and 320 8C deposition temperatures, respectively). Thermal cycling brought about massive changes in film stress. Representative stress-temperature data during thermal cycling are shown for Film 5 in Fig. 1. Near room temperature, reversible thermal stress develops; on further heating, irreversible increases in stress arise, leading to stress hysteresis (the magnitude of which is defined here as an increase in room temperature stress due to a thermal cycle). Stress always increased in the tensile (less compressive) direction, regardless of the absolute value of stress. Hence, this stress change process is not a relaxation of residual tensile or com-

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Film

1 2 3 4 5 6 7 8

Deposition temperature (8C)

Plasma frequency (MHz)

Substrate

320 320 320 320 120 120 120 120

13.56 13.56 0.46 0.46 13.56 13.56 0.46 0.46

InP Si InP Si InP Si InP Si

Deposition stress (MPa)

Plane-strain modulus (GPa)

Hardness (GPa)

CTE (25 8C) (10y6 8Cy1)

Failure conditions

340 y1010 y1160 y1440 130 y510 y600 y2280

152"7 185"5 192"16 214"5 140"5 180"4 133"9 194"8

16.5"0.3 18.2"0.3 21.3"1.2 23.8"0.6 14.5"0.5 16.0"0.2 15.6"0.8 19.4"0.4

2.8"0.1 1.4"0.1 1.9"0.3 1.5"0.1 3.4"0.2 3.1"0.3 2.0"0.3 3.0"0.2

500–630 )630 8C 200–500 500–630 500–630 )630 8C 300–400 )630 8C

Temperature (8C) 8C 8C 8C 8C 8C

Film stress (MPa)

Mechanism

)1330 )420 y1000–(y500) y830–0 1930–2750 )2050 y190–0 )1000

film cracking did not fail film delamination film delamination film cracking did not fail film delamination did not fail

M.P. Hughey, R.F. Cook / Thin Solid Films 460 (2004) 7–16

Table 1 Deposition conditions, properties and failure conditions for PECVD silicon nitride films

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Fig. 1. Representative stress-temperature data for a PECVD silicon nitride film (5, see Table 1) during thermal cycling. Illustrated are the beginning and end of thermal cycling, irreversible tensile stress development (curved arrow), the magnitude of stress hysteresis (straight arrow), and the linear reversible stress change caused by a thermal mismatch between film and substrate (dashed line).

pressive stress. This is illustrated more fully in Fig. 2, which shows the development of stress for films deposited at 320 8C and 13.56 MHz. Whereas Film 1 (on InP) was deposited in tension, Film 2 (on Si) was

Fig. 2. Stress-temperature data for thermal cycling of PECVD silicon nitride films deposited at 320 8C and 13.56 MHz on InP (Film 1) and Si (Film 2).

Fig. 3. Stress-temperature data for thermal cycling of PECVD silicon nitride films deposited at 120 8C and 0.46 MHz on InP (Film 7) and Si (Film 8).

deposited with a substantial compressive stress; both increased in stress at about the same rate. This stress hysteresis initiated at temperatures not much greater than the deposition temperature. In fact, a thermal cycle to 400 8C (only 80 8C above the deposition temperature) led to nearly 200 MPa stress development in both cases. On cycling above 500 8C, Film 1 cracked (discussed in Section 3.2), whereas Film 2 remained intact after thermal cycling to 630 8C, which induced over 1.1 GPa total stress change. Results for the films deposited at 120 8C and 0.46 MHz are shown in Fig. 3. After cycling to 300 8C, the stress state of Film 8 (on Si) had not changed, whereas Film 7 (on InP) had increased by nearly 400 MPa. Such a dramatic difference in cycling behavior was not observed for any other film pair. On cycling above 300 8C, Film 7 delaminated under compression (discussed in Section 3.2), whereas Film 8 survived all cycling and increased in stress by a total of over 2 GPa. (Stress-temperature data are shown only for complete thermal cycles in which the film did not fracture.) The stress hysteresis behavior is best examined in the manner of Fig. 4. In the top panel, the stress hysteresis is shown as a function of the difference between the peak cycle temperature, Tpeak, and the deposition temperature, Tdeposition. Films deposited at 120 8C required

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Fig. 4. A summary of the magnitude of stress hysteresis after each thermal cycle. Results are shown for stress hysteresis magnitude and film stress (of 120 8C films) as functions of the difference between peak cycle temperature and deposition temperature.

much greater temperature excursions from the deposition temperature to initiate stress hysteresis. However, in all cases, stress hysteresis occurred at relatively low temperatures (-400 8C). By examining the stress data for films that were able to survive at least three thermal cycles above the deposition temperature, a roughly linear increase in stress hysteresis is observed with increasing peak temperature, followed by a reduction in the hysteresis increase as the mechanism for stress change appears to near exhaustion (as in Film 6, for example). With the slight exceptions of the stress development in Films 4 and 7, the stress hysteresis appeared to be independent of the plasma frequency and substrate—and strongly dependent on deposition temperature. The magnitude of stress hysteresis is shown to be independent of film stress in the bottom panel, in which the film stress (sum of the deposition stress and the stress hysteresis) for the films deposited at 120 8C is plotted. The constant slope that was observed in the top panel is shown to remain constant over a range of film stress from greater than 2.2 GPa compression to greater than 1.4 GPa tension. As it is clear that the stress change is not a mechanical relaxation process, and the magnitude and sign of stress

do not affect the magnitude of hysteresis, the governing mechanism of stress change is not mechanical in nature. Instead, it is likely that a chemical change in the film produces lateral strains that cannot be accommodated due to the constraint of the stiff substrate, leading to stress development. As this stress change is always tensile, the chemical change must act to reduce the film volume. It is well known that many as-deposited PECVD films contain substantial amounts of bonded hydrogen (from the source gases) and that the amount of hydrogen in the film decreases with annealing. This loss of hydrogen, followed by network reforming, is surely responsible for the observed tensile stress development. Smith et al. w33x suggest that in high-quality silicon nitride films without Si–H bonds, film growth occurs by a condensation reaction that releases NH3; this occurs stress-free at the surface, but the uncompleted reaction continues in the bulk of the film to cause tensile stress. Those authors’ flash desorption experiments using mass spectrometry showed that NH3 indeed evolved from these types of films, although lesser amounts of H2 were also observed. Lu et al. w35x describe an additional hydrogen removal process for silicon nitride through H2, especially for films with Si–H bonds. Subsequent mass spectrometry studies during annealing indicated that mostly H2, but also some ammonia radicals, evolved from the films w36x. Both studies used infrared spectroscopy to show an increase in Si–N bonding with annealing. Furthermore, as ample evidence from hydrogenated amorphous silicon shows w37x, films with substantial amounts of Si–H bonds can produce H2 almost exclusively. Tensile stress generation in silicon oxide films has been explained by the reaction of neighboring silanols to form H2O and a network Si–O bond w6,38x, which is analogous to NH3 production in silicon nitride, with OH groups the chemical analogs of NH2 groups w39x. It is presumed here that the most likely process leading to tensile stress in common PECVD silicon nitride (containing Si–H and N–H bonds) is the simple chemical reaction of hydrogenbonded network-forming (Si–H or N–H) atoms to form molecular hydrogen and reformed network (Si–Si or Si–N), as shown pictorially in Fig. 5.

Fig. 5. An illustration demonstrating the likely chemical reactions responsible for stress development in silicon nitride with Si–H and N–H bonds.

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That this mechanism can account for the magnitude of the stress changes shown in Figs. 1–4 is easily demonstrated by considering the amount of hydrogen lost, the physical constraint of the substrate on the film, and the film biaxial modulus w6x. RBS and FReS measurements were conducted for the as-deposited high temperature films and all but the catastrophically fractured annealed films. RBS results showed that the film atomic density was approximately 1=1023 atomycm3; FReS results for the 320 8C deposition temperature films indicated that these films initially contained roughly 17– 20 at.% hydrogen, and all the annealed films contained 0.5–2 at.% hydrogen—meaning approximately 1.5– 2=1022 atomycm3 of hydrogen was lost. If a film were freestanding, the chemical reaction described above would lead to isotropic film shrinkage. However, the constraint of the substrate prevents any lateral shrinkage, and the volume strain is only partially accommodated by out-of-plane direction stress-free strain. In fact, the imposed lateral tensile strains, ´, are

´sDy3,

(3)

where D is the freestanding film relative volume reduction, or dilatation, equal to (VyV0)yV0, where V and V0 are the new and original film volumes, respectively. (This situation is identical to thermal stress development in which the film has a smaller CTE than the substrate and the CTE is isotropic.) The dilatation that would be imposed on a freestanding film is DsŽcyc0.Vy2,

(4)

where c and c0 are the new and original hydrogen bond concentrations (atomycm3 ), V is the reduction in freestanding film volume per chemical reaction (cm3), and the factor of 1y2 accounts for 2 hydrogen bond reactants per reaction. Finally, the induced biaxial film stress increase is calculated as sfsyE¯ f´syE¯ fŽcyc0.Vy6.

(5)

Taking Film 2 as an example (the film survived all thermal cycling and was nearly exhausted of hydrogen), the biaxial modulus of 230 GPa and total stress change of 1.11 GPa (a strain of 0.005) are used to calculate V ˚ 3, which is from Eq. (4); one obtains V;1.5–1.9 A equivalent to a ‘cube’ of reduced volume with an edge ˚ That this volume reduclength of approximately 1.2 A. tion is reasonable for the process depicted in Fig. 5 is further validation that such a mechanism is at least plausible. These large changes in stress resulted in large changes in the film CTE, as evidenced by a change in the room temperature stress-temperature slope after each thermal

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cycle. In this regard, the CTE is a better indicator than other mechanical properties of the structural changes in the film. The CTE values of Films 1 and 2 (deposited at 320 8C and 13.56 MHz) converge to nearly the same value (2.0–2.2 ppm Ky1) after repeated thermal cycling. Conversely, the CTE of Film 4 (deposited at 320 8C and 0.46 MHz on Si) decreased slightly. All films deposited at 120 8C exhibited a decrease in CTE to roughly 1.4 ppm Ky1. Finally, the CTE of Film 3 could not be measured, as the film fractured after being cycled directly to 630 8C. 3.2. Film fracture During thermal cycling five of the eight films fractured either by film cracking (caused by tensile film stress) or delamination (caused by compressive film stress). Fig. 6 displays optical micrographs of four of the fractured film surfaces (Films 1, 3, 4 and 7); Film 5 cracked in a very similar manner to Film 1. Film 1 was observed to crack, with cracks oriented primarily in one direction—indicative of a biaxial stress that was larger in one direction (perpendicular to cracks) than the other. This was also evident in the large uncertainties in the calculation of average film stress. Despite the very large film stress and availability of a free edge created by cracking, little delamination occurred, indicating a large interfacial fracture resistance. Films 2 and 3 were deposited under identical conditions and both delaminated under compression at elevated temperatures, but circular delaminations in Film 4 (on Si) were of very uniform size, whereas the circular delaminations in Film 3 (on InP) were much smaller and of non-uniform size, and numerous defects combined to form very large buckled regions. The large disparity in the number of delaminated regions between the two films suggests that the density of defects that allow such delamination was much greater on the film-InP interface. Finally, Film 7 fractured in much the same way as Film 3, however, the circular regions were much smaller still, and they combined to form large buckled regions with more linear features, not circular or ring-shaped. Perhaps more informative than the visual observation of the fractured films are the conditions under which the films fractured (or survived). Table 1 lists the temperature and film stress at which the films fractured (or the most severe conditions survived by the films) along with brief fracture descriptions. Temperatures and stresses are expressed in ranges, as the fracture conditions were not known precisely (it is not possible to visually inspect the films at all times during thermal cycling). It is evident that films survived much larger values of tension (over 2 GPa in some cases) than compression, even at temperatures as high as 630 8C. This is not necessarily surprising, as the fracture associated with tension is entirely within the film, whereas

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Fig. 6. Optical micrographs of the fractured film surfaces of (a) Film 1, (b) Film 3, (c) Film 4, and (d) Film 7.

compressive fracture occurs at the film-substrate interface. Also, the magnitudes of compressive stress under which films failed were smaller than the as-deposited states, indicating that it was thermal activation, not an increase in stress, which initiated compressive fracture. No tensile films fractured in a corresponding manner (high temperature, small stress); thus it is not clear how important thermal activation was for film cracking. Differences in fracture phenomena for films on different substrates were observed. Films on InP cracked more easily, apparently because they were always deposited in less compression, forcing the films into larger values of tension after a given thermal cycle. However, films on Si were not more likely to delaminate under compression despite larger values of compression. This is most evident on comparison of Films 7 and 8: Film 7 delaminated under near zero compression at approximately 400 8C, whereas Film 8 did not delaminate while under 400 MPa to 2 GPa compression at 400–500 8C. Clearly, the interface plays a key role for the films that failed in compression. Naturally, the interfacial fracture resistance should be different for identical films on different substrates. However, the different delamination

fracture morphologies are perhaps better explained by the presence of stable native oxide (on Si) vs. an unstable oxide (on InP) on the ‘bare’ substrates. It is likely that the stable native SiO2 oxide provided fewer defects at which delamination events could nucleate. This possibly provided a greater resistance to delamination and hindered the coalescence of multiple fracture events into massive buckled regions. The diverse fracture phenomena and conditions did not have a clear dependence on deposition conditions. The effect of plasma frequency on fracture was not evident in three of the four film pairs deposited at the same temperature on the same substrate, as they either survived all thermal cycling or did not fail in the same manner because the deposition stress dictated either a tensile or compressive failure. However, films deposited at 320 8C on Si do offer one potential comparison. Films 2 and 4 were deposited in nearly the same level of compression, but Film 4 delaminated and Film 2 did not. It is possible, then, that the low plasma frequency used in the deposition of Film 4 reduced interfacial defects or provided a greater interfacial fracture resistance. However, Film 2 exhibited tensile stress above

M.P. Hughey, R.F. Cook / Thin Solid Films 460 (2004) 7–16

500 8C, whereas Film 4 was still in compression at these temperatures. It is more likely that the difference in stress was responsible for the different fracture response. Direct comparisons of film pairs deposited at the same frequency on the same substrate are similarly hampered, but two film pairs provide some understanding. Films 1 and 5 (deposited at 13.56 MHz on InP) both cracked, however, Film 5 (deposited at 120 8C) survived a much larger film stress before fracturing. Whether the better performance of Film 5 is due to an increase in fracture toughness or the cracking threshold is unclear, but either is beneficial. Also, Films 4 and 8 (deposited at 0.46 MHz on Si) exhibited different fracture behavior, as Film 4 (deposited at 320 8C) delaminated and Film 8 (deposited at 120 8C) did not fracture, despite both films being deposited in a substantial amount of compression. However, because the stress hysteresis of the Film 8 was much greater at elevated temperatures, Film 4 was more compressive at high temperature, leading to its fracture. The estimated stress at which failure occurred allows fracture resistance approximations to be made for the films that cracked. For the through-film cracking and subsequent delamination observed in Films 1 and 5, the appropriate equation for the mode-I and mode-II stressintensity factors is w40x KIsKIIscsftf1y2

(6)

where c is a geometry term. For film cracking, the mechanical energy release rate, or crack driving force, is GsK2I yE,

(7)

and c;1.41 w40x for the approximately moduli-matched film-substrate system w41x. At equilibrium, Gs2g, where 2g is the surface energy density in air and is the film fracture resistance. By using the approximate stress values at fracture, Eqs. (6) and (7) give 2g;5.9 Jym2 for Film 1 and 2g;14.2 Jym2 for Film 5. (Corresponding toughness values are 0.9 MPa m1y2 for Film1 and 1.4 MPa m1y2 for Film 5.) Because some delamination initiated from the free edge created by cracking (see insert on Fig. 6a) and stabilized, bounds can be placed on the interfacial fracture resistance, R, where GsR at equilbrium. The delamination crack driving force is w40x GsŽ1yE*f q1yE*s .ŽK2I qK2II.y2.

(8)

The fracture resistance was clearly less than the initial driving force, for which cs0.71, and was clearly greater than the long delamination driving force, for which cs 0.43 for KI and 0.56 for KII w40x. After taking an average

15

value of E*s over all in-plane directions, R is calculated to be between 2.1 and 4.1 Jym2 for Film 1 and between 4.7 and 9.4 Jym2 for Film 5. Despite the additional fracture mode involved in delamination, the interfacial fracture resistance is still smaller than the film fracture resistance. It should be noted that all of these calculations neglect the effect of high temperature during fracture of these films. 4. Conclusions Massive stress changes (up to 2 GPa) are possible in hydrogen-containing PECVD silicon nitride on thermal cycling. Irreversible stress development initiated at moderate temperatures (below 400 8C) for all films, and other deposition parameters did not appear to affect the magnitude of hysteresis. This stress hysteresis was observed to be independent of film stress, indicating that a chemical change within the film was the driving force, not mechanical relaxation. A general mechanism for stress change, proposed by others, has been adopted here and was shown to account for the observed magnitude of stress change. The hydrogen release from the films necessarily induces a change in film structure, which was manifested in a change in film CTE: the film CTE typically decreased on thermal cycling, but not for all films. The influence of the choice of substrate on the deposition stress and properties of these amorphous films is not clear at this time. However, the influence of the substrate on film fracture was dramatic and is understandable in some cases. Films on InP had larger deposition stresses, causing them to crack more easily during the irreversible stress development that accompanied thermal cycling. Additionally, an increased defect density was evident on film-InP interfaces, leading to varied fracture morphologies on this substrate and a smaller resistance to compressive fracture. Fracture properties appeared to be largely independent of most deposition parameters, although deposition with low frequency plasma may have resulted in a larger film toughness. Fracture properties were strongly dependent on the magnitude of stress, as films were able to survive massive magnitudes of tension at high temperatures, but frequently delaminated under relatively small magnitudes of compression and at modest temperatures. If high temperature (here, 630 8C) processing or use is required of a PECVD film, then high temperature, high frequency deposition is the optimal choice. High temperature is preferred because less total stress hysteresis occurs. Both films deposited at 13.56 MHz on Si survived all cycling, and while the same films deposited on InP fractured, better design could reduce cracking (for example, the films could be deposited in slightly greater compression). The compressive fractures of Films 3, 4 and 7 do not necessarily rule out films

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deposited at 0.46 MHz from use. The very large compressive stresses afforded by this deposition choice are desired because they allow for greater total stress change. To avoid the compressive fractures observed here, the rate of thermal cycling can be manipulated such that slower cycling results in a conversion from compressive to tensile stress at relatively low temperatures (300–400 8C, depending on the film); then the rate of thermal cycling can be increased, and the films can ‘safely’ develop tensile stress. Results shown in this work provide further evidence for the generality of phenomena observed in different PECVD dielectric materials—primarily, stress increases irreversibly on thermal cycling and is caused by a reduction in the amount of bonded hydrogen. The increase in stress is observed here to be completely independent of film stress. As such, any future development of quantitative models of stress-temperature behavior for dielectric films, analogous to those for metal films, should be based on chemical changes in the film, with stress ‘along for the ride’ as a mechanical consequence. References w1x M.D. Thouless, J. Gupta, J.M.E. Harper, J. Mater. Res. 8 (1993) 1845. w2x R.P. Vinci, E.M. Zielinski, J.C. Bravman, Thin Solid Films 262 (1995) 142. w3x R.-M. Keller, S.P. Baker, E. Arzt, Acta Mater. 47 (1999) 415. w4x W. Zeiger, W. Bruckner, ¨ J. Schumann, W. Pitschke, H. Worch, Thin Solid Films 370 (2000) 315. w5x M.J. Kobrinsky, C.V. Thompson, M.E. Gross, J. Appl. Phys. 89 (2001) 91. w6x J. Thurn, R.F. Cook, J. Appl. Phys. 91 (2002) 1988. w7x X. Zhang, K.-S. Chen, R. Ghodssi, A.A. Ayon, ´ S.M. Spearing, Sens. Actuat. A 91 (2001) 373. w8x R.C. Budhani, R.F. Bunshah, P.A. Flinn, Appl. Phys. Lett. 52 (1988) 284. w9x C.W. Lim, C.C. Toh, J.K. Fu, S.K. Lahiri, J. Mater. Sci. Lett. 15 (1996) 2177. w10x M.A. El Khakani, M. Chaker, A. Jean, S. Boily, H. Pepin, ´ J.C. Kieffer, S.C. Gujrathi, J. Appl. Phys. 74 (1993) 2834. w11x J. Thurn, R.F. Cook, M. Kamarajugadda, L.C. Stearns, in: C.S. Ozkan, L.B. Freund, R.C. Cammarata, H. Gao (Eds.), Thin Films: Stresses and Mechanical Properties IX, Boston, USA, November 26–30, 2001, Materials Research Society Symposium Proceedings 695 2002, pp. L3.9.

w12x D.R. Harding, L.U.T. Ogbuji, M.J. Freeman, J. Appl. Phys. 78 (1995) 1673. w13x R.A. Gardner, P.J. Peterson, T.N. Kennedy, J. Vac. Sci. Technol. 14 (1977) 1139. w14x J.-H. Jou, L. Hsu, S. Yeh, T. Shyy, Thin Solid Films 201 (1991) 69. w15x R. Thielsch, A. Gatto, N. Kaiser, Appl. Opt. 41 (2002) 3211. w16x K.-S. Chen, X. Zhang, S.-Y. Lin, Thin Solid Films 434 (2003) 190. w17x C. Gorecki, Opt. Lasers Eng. 33 (2000) 15. w18x G.G. Stoney, Proc. R. Soc. Lond. A 82 (1909) 172. w19x W.A. Brantley, J. Appl. Phys. 44 (1973) 534. w20x S. Adachi, Physical Properties of III-V Semiconductor Compounds: InP, InAs, GaAs, GaP, InGaAs, and InGaAsP, John Wiley and Sons, New York, 1992, p. 23. w21x D.N. Nichols, D.S. Rimai, R.J. Sladek, Solid State Commun. 36 (1980) 667. w22x A.K. Bhattacharya, W.D. Nix, Int. J. Solids Struct. 24 (1988) 1287. w23x R.B. Roberts, J. Phys. D 14 (1981) L163. w24x R. Bisaro, P. Merenda, T.P. Pearsall, Appl. Phys. Lett. 34 (1979) 100. w25x G. Himsolt, H. Knoch, H. Huebner, F.W. Kleinlein, J. Am. Ceram. Soc. 62 (1979) 29. w26x T.F. Retajczyk Jr, A.K. Sinha, Thin Solid Films 70 (1980) 241. w27x M. Maeda, K. Ikeda, J. Appl. Phys. 83 (1998) 3865. w28x I. Tanaka, G. Pezzotti, T. Okamoto, Y. Miyamoto, M. Koizumi, J. Am. Ceram. Soc. 72 (1989) 1656. w29x O. Unal, J.J. Petrovic, T.E. Mitchell, J. Mater. Res. 8 (1993) 626. w30x R.N. Katz, Science 208 (1980) 841. w31x D.L. Smith, in: G. Lucovsky, D.E. Ibbotson, D.W. Hess (Eds.), Characterization of Plasma-Enhanced CVD Processes, Boston, USA, November 27–28, 1989, Materials Research Society Symposium Proceedings 165 1990, pp. 69. w32x R.H. Bruce, J. Appl. Phys. 52 (1981) 7064. w33x D.L. Smith, A.S. Alimonda, C.-C. Chen, S.E. Ready, B. Wacker, J. Electrochem. Soc. 137 (1990) 614. w34x A.C. Adams, F.B. Alexander, C.D. Capio, T.E. Smith, J. Electrochem. Soc. 128 (1981) 1545. w35x Z. Lu, P. Santos-Filho, G. Stevens, M.J. Williams, G. Lucovsky, J. Vac. Sci. Technol. A 13 (1995) 607. w36x P. Santos-Filho, G. Stevens, Z. Lu, K. Koh, G. Lucovsky, in: J.S. Im, B. Park, A.L. Greer, G.B. Stephenson (Eds.), Thermodynamics and Kinetics of Phase Transformations, Boston, USA, November 27–30, 1995, Materials Research Society Symposium Proceedings 398 1996, pp. 345. w37x W. Beyer, H. Wagner, J. Appl. Phys. 53 (1982) 8745. w38x M.S. Haque, H.A. Naseem, W.D. Brown, J. Appl. Phys. 82 (1997) 2922. w39x D.L. Smith, J. Vac. Sci. Technol. A 11 (1993) 1843. w40x M.D. Thouless, J. Vac. Sci. Technol. A 9 (1991) 2510. w41x J.L. Beuth Jr, Int. J. Solids Struct. 29 (1992) 1657.